Correlation of microstructure with tensile, creep and corrosion ...

Correlation of microstructure with tensile, creep and

corrosion behaviour of AZ91 Mg alloy fabricated by three

different casting techniques

A thesis submitted in partial fulfilment of the

requirements for the degree of

Master of Technology (By Research)

in

Metallurgical and Materials Engineering

by

Hrishikesh Shastri

(Roll No.: 612MM3011)

Department of Metallurgical and Materials Engineering

National Institute of Technology Rourkela, Odisha

2015

Correlation of microstructure with tensile, creep and

corrosion behaviour of AZ91 Mg alloy fabricated by three

different casting techniques

A thesis submitted in partial fulfilment of the

requirements for the degree of

Master of Technology (By Research)

in

Metallurgical and Materials Engineering

by

Hrishikesh Shastri

(Roll No.: 612MM3011)

Under the supervision of

Prof. A. K. Mondal and Prof. K. Dutta


National Institute of Technology Rourkela, Odisha

2015

Contents

Certificate .................................................................................................................................. i

Acknowledgements .................................................................................................................. ii

Abstract .................................................................................................................................... iii

List of figures ........................................................................................................................... iv

Tables ....................................................................................................................................... vi

Chapter 1. Introduction ....................................................................................................... 1-3

1.1 Introduction ...................................................................................................................... 1

1.2 Objective of the present thesis ......................................................................................... 2

1.3 Outline of the thesis .......................................................................................................... 3

Chapter 2. Literature review ............................................................................................. 4-16

2.1 Mg and its alloys .............................................................................................................. 4

2.1.1 Pure Mg ..................................................................................................................... 4

2.1.2 AZ91 alloy ................................................................................................................. 6

2.1.3 Applications of Mg alloys ......................................................................................... 6

2.2 Tensile test ........................................................................................................................ 8

2.3 Creep ................................................................................................................................ 8

2.4 Corrosion ........................................................................................................................ 11

2.5 Casting techniques ......................................................................................................... 12

2.5.1 Ingot casting (IC)..................................................................................................... 12

Contents

2.5.2 Squeeze-casting (SC) .............................................................................................. 12

2.5.3 High pressure die-casting (HPDC) .......................................................................... 12

2.6 Tensile and creep behaviour of Mg alloys in different casting conditions .................... 14

2.7 Corrosion behaviour of Mg alloys in different casting conditions ................................. 16

Chapter 3. Experimental procedure ............................................................................... 17-21

3.1 Fabrication of the alloys ................................................................................................. 17

3.2 Specimen preparation ..................................................................................................... 19

3.3 Characterization of the alloys ......................................................................................... 19

3.3.1 X-ray diffraction (XRD) .......................................................................................... 19

3.3.2 Optical microscopy (OM) ....................................................................................... 19

3.3.3 Scanning electron microscopy (SEM) ..................................................................... 20

3.3.4 Mechanical characterization .................................................................................... 20

3.3.4.1 Tensile test .......................................................................................................... 20

3.3.4.2 Creep test ........................................................................................................... 20

3.3.5 Corrosion test ......................................................................................................... 20

Chapter 4. Microstructural characterization ................................................................. 22-29

4.1 As-cast microstructure .................................................................................................... 22

Chapter 5. Mechanical characterization......................................................................... 30-40

5.1 Tensile behaviour ........................................................................................................... 30

5.2 Creep behaviour ............................................................................................................. 37

Contents

Chapter 6. Corrosion behaviour ...................................................................................... 41-56

6.1 Corrosion response ......................................................................................................... 41

6.1.1 Open circuit potential (OCP) ................................................................................... 41

6.1.2 Electrochemical impedance spectroscopy (EIS) ..................................................... 43

6.1.3 Potentiodynamic polarization scan .......................................................................... 49

6.2 Microstructural observation after corrosion tests ........................................................... 51

Chapter 7. Conclusions ..................................................................................................... 57-58

References .......................................................................................................................... 59-61

Bio-data ................................................................................................................................... 62

i


National Institute of Technology Rourkela

Odisha-769008, India

CERTIFICATE

This is to certify that the thesis entitled “Correlation of microstructure with tensile, creep and

corrosion behaviour of AZ91 Mg alloy fabricated by three different casting techniques”

submitted by Hrishikesh Shastri (Roll No.: 612MM3011) in partial fulfilment of the

requirements for the award of Master of Technology (By Research) in Metallurgical and

Materials Engineering at National Institute of Technology, Rourkela is a bonafide research

work carried out by him under our supervision and guidance and is worthy of consideration for

the award of the degree. The candidate has fulfilled all prescribed requirements for the thesis.

To the best of our knowledge, the matter embedded in the thesis is based on candidate’s own

work and has not been submitted to any other university / institute for the award of any degree

or diploma.

Prof. Ashok Kumar Mondal Prof. Krishna Dutta

Principal Supervisor Co-Supervisor

Date:

Place:

ii

Acknowledgements

My sincere thanks to the present and past HOD, Department of Metallurgical and Materials

Engineering, National Institute of Technology Rourkela for providing the necessary facilities

for carrying out this work.

My heartfelt thanks goes to my principal supervisor Prof. A. K. Mondal for believing in me.

Without him this work would have not been possible. The discussion with him were

enlightening. I am also highly thankful to my co-supervisor Prof. K. Dutta for his constant

support and encouragement.

My sincere gratitude to MSC members Prof. S. C. Mishra, Prof. A. Basu, Prof. M. Masanta

and Prof. A. Mitra for their guidance throughout this work.

I am also thankful to my labmates Anil and Pranav for their support. I would like to thank all

the laboratory staffs of the Dept. of MM, NIT Rourkela for their constant practical assistance.

I am also grateful to the Biju Patnaik Central Library for providing me the required literatures

and books for carrying out this research work.

Thanks to all help and company I received from the Dept. of Materials Engineering, Indian

Institute of Science, Bangalore. It would have not been possible without the support extended

by the NRC-M. Prof. S. Kumar and Prof. S. Suwas have always been so helpful. My special

thanks goes to Chandra Shekhar P. and his labmates for their commendable help with the work.

I also thank Dr. H. Dieringa for helping me to carry out the creep tests.

Last but not the least, I feel grateful to have such a supportive family and friends.

Hrishikesh Shastri

iii

Abstract

In the present investigation an attempt has been made to correlate microstructure with tensile,

creep and corrosion behaviour of the AZ91 Mg alloy produced by three different casting

techniques. All the as-cast alloys consist of primary Mg (α-Mg) and β-Mg17Al12 phases. The

volume fraction of the β-Mg17Al12 phase is the highest in the ingot casting (IC), intermediate

in the squeeze-casting (SC) and the lowest in the high pressure die-casting (HPDC). The tensile

tests are performed at ambient, 150 and 200°C temperatures. The best tensile properties are

exhibited by the SC alloy at all the temperatures employed in the present investigation owing

to the presence of negligible porosity and relatively finer grain size. Fracture surfaces of the

broken tensile specimens reveal quasi-cleavage fracture. The creep tests are performed in

compression at a stress of 70 MPa and temperature of 150ºC. The best and worst creep

resistance are exhibited by the SC and HPDC alloys with the IC alloy exhibiting the

intermediate creep resistance. Poor creep resistance of the HPDC alloy is attributed to the

presence of relatively higher amount of porosity that allows easy crack initiation and growth

during creep tests. On the contrary, presence of negligible porosity and continuous network of

eutectic phase contributes to the superior creep resistance of the SC alloy. Electrochemical

corrosion tests of all the alloys are carried out at ambient temperature in 0.5 wt.% NaCl aqueous

solution with neutral pH. The best and worst corrosion resistance are exhibited by the HPDC

and IC alloys. The superior corrosion resistance of the HPDC alloy is attributed to the relatively

finer grain size and higher percentage of Al dissolved in α-Mg phase. Thus, the SC alloy

exhibits the best tensile and creep properties with intermediate corrosion resistance among the

alloys employed in the present investigation.

Keywords: AZ91 magnesium alloy; Casting technique; Microstructure; Tensile; Creep;

Corrosion

iv

List of figures

Fig. No. Caption Page No.

2.1 Schematic description of HCP crystal structure and major planes

of Mg.

5

2.2 Mg-Al phase diagram. 7

2.3 Schematic representation of an engineering stress-strain curve. 9

2.4 Schematic representation of a typical creep curve. 9

2.5 Schematic diagrams of various operations of direct SC process:

(a) pouring melt into die cavity, (b) solidification under pressure

and (c) removal of casting.

13

2.6 Schematic diagram of cold chamber HPDC. 13

3.1 Actual photograph of the SC setup employed to fabricate the IC

and SC alloys.

18

4.1 XRD patterns obtained from the AZ91 alloy in three different

casting conditions.

23

4.2 Optical micrographs of the (a) IC, (b) SC and (c) HPDC alloys. 24

4.3 SEM micrographs of the (a) IC, (b) SC and (c) HPDC alloys. 26

5.1 Stress vs. strain plots of the SC alloy tested at ambient, 150°C

and 200°C temperature.

31

5.2 Comparison of tensile properties of the IC, SC and HPDC alloys

at (a) ambient, (b) 150°C and (c) 200°C temperature.

32

5.3 Fracture surfaces of tensile specimens of the (a) IC, (b) SC and

(c) HPDC alloys tested at ambient temperature.

36

List of figures

v

5.4 (a) Strain vs. time plots for the IC, SC and HPDC alloys creep

tested at 70 MPa stress and temperature of 150°C for 50 h; (b)

Strain rate vs. time plots and (c) Calculation of strain rates by

linear fit to the plots shown in (a).

38

5.5 SEM micrograph of the HPDC alloy creep tested at 70 MPa

stress and temperature of 150°C for 50 h.

40

6.1 Variation of OCP with time of the HPDC, SC and IC alloys. 42

6.2 Comparison of OCP of the HPDC, SC and IC alloys. 42

6.3 Nyquist plots of the (a) HPDC, (b) SC and (c) IC alloys. 45

6.4 Nyquist plots of all the alloys at (a) 0 h, (b) 10 h and (c) magnified

view of (b).

46

6.5 Bode impedance plots of the (a) HPDC, (b) SC and (c) IC alloys. 47

6.6 Bode phase angle plots of the (a) HPDC, (b) SC and (c) IC alloys. 48

6.7 Potentiodynamic polarization plots for all the alloys. 50

6.8 Comparison of corrosion rate obtained from Tafel extrapolation. 50

6.9 SEM micrographs of the surfaces corresponding to the (a)

HPDC, (b) SC and (c) IC alloys before removal of corrosion

products.

52

6.10 XRD pattern obtained from the corroded surface of the SC alloy. 53

6.11 FTIR pattern of corrosion product obtained from the corroded

surface of the SC alloy.

53

6.12 SEM micrographs of the (a) HPDC, (b) SC and (c) IC alloys after

removal of corrosion products.

55

6.13 SEM micrograph of cross-section of the corroded surface of the

IC alloy.

56

vi

Tables

Table No. Caption Page No.

2.1 Applications of Mg alloys in automobiles. 6

4.1 Quantitative analysis of the microstructural features of all the

alloys.

28

5.1 Comparison of the calculated and experimentally determined

YS of all the alloys.

33

6.1 Values of various parameters obtained from the

potentiodynamic polarization plots.

49

1

Chapter 1. Introduction

1.1 Introduction

Environmental interest and exhaustible reserves of fossil fuel demands for weight-saving to

attain improved fuel efficiency in the automotive industry. In this regard, magnesium (Mg)

alloys possessing high specific strength are the most promising structural material for

applications in aerospace and automotive industries. Mg alloys are primarily used in as-cast

condition to fabricate intricate components and the casting conditions or parameters (i.e.,

solidification rate, volume fraction and morphology of intermetallic phases, grain size,

porosity, degree of solute supersaturation and so on) decides the achieved properties.

Among the various casting techniques available, ingot casting (IC) and high pressure die-

casting (HPDC) are widely used in automotive applications. IC is a conventional casting

process where the molten metal solidifies under the influence of gravity. In HPDC, molten

metal is injected at high speed into the mould and solidification occurs under high pressure.

The solidification in HPDC is very rapid (varies from milli-seconds to a few seconds) and gives

a net-shaped component. Recently, squeeze-casting (SC) gained popularity for the superior

properties exhibited by the castings produced following this route. In SC, the rate of molten

metal pouring is relatively slower, continuous and solidification takes place under high pressure

that ensures negligible porosity in the casting. The governing reason for low porosity in the

SC as compared to HPDC is that the former takes longer time for solidification (about a half-

second to three seconds) and lower velocity of die filling as compared to the latter ensuring

less air-entrapment.


2

Various Mg alloys with different alloying additions were produced. However, the AZ91 (Mg-

9Al-0.7Zn-0.2Mn) (wt.%) alloy is a cheaper alternative with decent properties and continues

to be a workhorse alloy. A review of the existing literature (presented in chapter 2) evidenced

that the tensile, creep and corrosion behaviour of Mg alloys differ significantly with cooling

rates and casting routes employed to fabricate the alloys. To the best of our knowledge, no

literature exists that correlates microstructure, mechanical properties and corrosion behaviour

of the AZ91 alloy with three different casting techniques i.e., IC, SC and HPDC side-by-side.

Consequently, it is worth correlating microstructure, mechanical properties and corrosion

behaviour of the AZ91 alloy fabricated by three different casting techniques in order to

ascertain the most suitable casting route for large-scale production.

1.2 Objective of the present thesis

The objective of the present investigation is to correlate microstructure with tensile, creep and

corrosion behaviour of the AZ91 Mg alloy fabricated by three different casting techniques.

The salient objectives of this investigation are as follows:

i. Fabrication of the AZ91 alloy by three different casting routes i.e., IC, SC and HPDC.

ii. Characterization of microstructures of all the fabricated alloys.

iii. Characterization of ambient and elevated temperature tensile properties as well as creep

behaviour of the alloys.

iv. Observation of microstructures of tensile and creep tested specimens in order to

understand the underlying deformation mechanisms.

v. Investigation of corrosion behaviour of all the fabricated alloys.


3

1.3 Outline of the thesis

The present thesis contains seven chapters and the contents of each chapter are mentioned as

follows:

Chapter 1 discusses the relevance of the problem.

Chapter 2 reviews the literature most relevant to the present investigation.

Chapter 3 describes the experimental procedures followed.

Chapter 4 presents the results and discussion on microstructural characterization of the as-cast

alloys.

Chapter 5 includes the results and discussion on tensile and creep behaviour of the alloys.

Chapter 6 deals with the investigation of corrosion behaviour of the alloys.

Chapter 7 concludes the present investigation.

The references cited throughout the thesis are listed in the reference section.

4

Chapter 2. Literature review

This chapter briefly discusses the literature most relevant to the present investigation. Section

2.1 is a brief introduction about Mg and its alloys. Section 2.2 and 2.3 provide summary about

tensile and creep tests. Section 2.4 includes fundamentals of corrosion. The details about

different casting techniques have been discussed in section 2.5. Literature covering mechanical

properties and corrosion behaviour of Mg alloys in different casting conditions is reviewed in

sections 2.6 and 2.7, respectively.

2.1 Mg and its alloys

2.1.1 Pure Mg

Mg (density: 1.74 gm/cm3 and 35% lesser in weight than Al) is the lightest structural metal and

therefore, has the abundant possibilities for applications in automobile and aerospace sectors.

In addition, it has high specific strength, good castablity, good damping capacity, good

machinability, good weldability in controlled environment and considerably developed

corrosion resistance with high purity Mg [1]. Mg has HCP crystal structure as shown in Fig.

2.1 with lattice parameters a 3.18 Å and c 5.19 Å, which is nearly identical to the ideal c/a

ratio of 1.62354 (at 25°C) [2]. Mg has limited ductility at ambient temperature owing to three

active slip systems, however, at elevated temperature prismatic and pyramidal slip planes also

get activated resulting in higher ductility.


5

Fig. 2.1. Schematic description of HCP crystal structure and major planes of Mg [2].


6

2.1.2 AZ91 alloy

AZ91 Mg alloy (Mg-9Al-0.7Zn-0.2Mn) (wt.%) is the most commonly used alloy in as-cast

condition. It is economical and delivers excellent castablity and excellent ambient temperature

mechanical properties. Al as an alloying element in AZ91 improves its castablity by widening

the freezing range. It enhances strengthening by forming solid solution and by precipitation of

the intermetallic phase (Mg17Al12) at low temperatures (<120°C). The Mg-Al phase diagram

[3] is shown in Fig. 2.2. Zn also increases room-temperature strength. It counteracts the

corrosive effect of Ni and Fe impurities. Small amount of Mn is added to enhance saltwater

resistance by forming intermetallic compounds with impurities like Fe [4].

2.1.3 Applications of Mg alloys

Some automotive applications of Mg alloys are shown in Table 2.2.

Table 2.1. Applications of Mg alloys in automobiles [5].

System Product

Interior Instrument panel, seat frame, steering wheel

Body Door frame, roof frame

Chassis Wheel, brake pedal bracket

Powertrain Manual transmission case, oil pan


7

Fig. 2.2. Mg-Al phase diagram [3].


8

2.2 Tensile test

Tensile tests are performed for different purposes such as material selection for engineering

applications, quality inspection and novel material development. The data obtained are usually

plotted as stress vs. strain. A typical stress-strain curve is shown in Fig. 2.3. In the elastic

region, stress is linearly proportional to strain. When the stress exceeds a value corresponding

to the yield strength, the specimen undergoes gross plastic deformation. If the load is

subsequently reduced to zero, the specimen will remain permanently deformed. For

engineering applications, the elastic limit is defined by the offset yield strength (generally at

strain = 0.002). Initially, a balance exists between strain hardening and reduction in the area,

and the engineering stress increases with increasing strain. Ultimately, the load reaches a

maximum value after which the decrease in specimen’s cross-sectional area is greater than the

increase in deformation load due to strain hardening and the specimen begins to neck locally.

Therefore, the actual load required to deform the specimen decreases and it finally fractures

[6].

2.3 Creep

Creep is defined as the time-dependent plastic deformation of materials when subjected to

constant load or stress at constant temperature ( 0.4Tm, where Tm is the absolute melting

temperature). It is an undesirable phenomenon which often limits the lifetime of a component

in service. Fig. 2.4 shows a typical creep curve of metal at constant stress and constant

temperature. Upon load application, there is an instantaneous elastic deformation. The resulting

creep curve consists of three regions. Primary or transient, secondary and tertiary creep.

Primary creep is characterised by decrease in slope of the curve with increase in time

suggesting an increase in creep resistance. In secondary stage, also known as steady-state creep,

a balance between the competing processes of strain hardening and recovery takes place and


9

Fig. 2.3. Schematic representation of an engineering stress-strain curve [6].

Fig. 2.4. Schematic representation of a typical creep curve [7].


10

the plot becomes linear. In tertiary stage, creep rate increases and material fails. This failure

generally results from grain boundary separation, formation of internal cracks, cavities, voids

and so on [7].

Creep deformation is generally classified as:

1. Dislocation creep

2. Diffusional creep

3. Grain boundary sliding

Dislocation creep

The movement of dislocations during deformation overcoming the obstacles by either glide

and/or climb is called dislocation creep. This occurs by the help of thermally assisted

mechanisms involving the diffusion of interstitials or vacancies. The glide along the slip planes

are referred to as glide-controlled creep, whereas, the climbing up of dislocations to parallel

slip plane across hindrances is called climb controlled creep. At high temperatures, the key

mechanism for overcoming obstacles is dislocation climb and the creep rate is not dependent

on grain size. The creep rate is given as,

ε�̇� =A𝐷𝑣𝐺𝑏

𝑘𝑇(σ

𝐺)𝑛 (1)

where, A and n are the material constants, Dv is diffusion coefficient, G is shear modulus, b is

Burgers vector, k is Boltzmann’s constant and T is absolute temperature [9]. The value of n

varies from 3 to 8. The activation energy for climb-controlled creep is the same as the activation

energy for self-diffusion [8].


11

Diffusional creep

It involves motion of vacancies and atoms through a metal crystal under the influence of

applied stress. It occurs at relatively low stresses, σ/G < 10-4 [9]. If creep occurs by diffusion

through lattice, it is called Nabarro-Herring creep, whereas, grain-boundary diffusion

predominates in Coble creep.

Grain boundary sliding

It involves sliding of neighbouring grains with respect to the grain boundary separating them.

At temperatures above 0.5Tm, the grain boundaries have less viscosity and acts as a viscous

liquid allowing grain boundary sliding.

2.4 Corrosion

Corrosion is defined as the deterioration of materials due to reaction with environment. The

corrosion products may dissolve or remain in solid form. The corrosion tests are useful for

prediction of corrosion behaviour, material selection and study of corrosion mechanism. As

corrosion is an electrochemical process, therefore, electrochemical techniques are used for the

study of corrosion process. Electrochemical techniques are popular because they are fast,

accurate and versatile. The tests are carried out in a potentiostat which is an electronic

instrument that controls the potential difference between a working electrode (WE) and a

reference electrode (RE) present in a single electrochemical cell by introducing current into the

cell through an auxiliary or counter electrode. The potential of the WE relative to RE when no

current or potential is applied to the cell is called as open circuit potential (OCP) or corrosion

potential (Ecorr). At Ecorr, rate of oxidation is equal to the rate of reduction. Electrochemical

spectrum is obtained by the linear potential sweep technique also known as potentiodynamic

polarization scan. It offers significant information on the corrosion rate, passivity, as well as


12

the relative cathodic and anodic behaviour of an electrochemical system. Detailed examination

of reaction mechanisms including passive layer investigation can be carried out by

electrochemical impedance spectroscopy (EIS) method. EIS uses a scan of test frequencies

normally from a high to a low frequency to analyse resistive and capacitive characteristics of a

corrosion cell. Data is presented in either Nyquist, Bode Phase or Bode Impedance formats.

2.5 Casting techniques

2.5.1 Ingot casting (IC)

It is a type of gravity casting similar to conventional sand casting method. This process uses

metal mould. Compared to sand casting, it has advantages including superior surface finish,

greater dimensional control and better mechanical properties owing to faster solidification.

However, there are limitations on casting intricate components through IC.

2.5.2 Squeeze-casting (SC)

There are two types of SC processes: direct and indirect. In both processes, molten metal is

introduced in mould with minimum turbulence and solidification takes place under high

pressure (typically above 100 MPa) within closed dies. Direct SC is also called liquid metal

forging in which molten metal is poured into a preheated, lubricated die and the pressure is

applied till solidification takes place [10]. A schematic of SC setup is shown in Fig. 2.5.

2.5.3 High pressure die-casting (HPDC)

Intricate components of light metals can be easily cast by HPDC. The excellent castablity of

Mg alloys are fully utilized in this process. There are mainly two types of HPDC processes:

hot chamber HPDC and cold chamber HPDC. In the hot chamber HPDC method, the molten

metal is kept in a steel crucible, under a protective atmosphere, whereas, protective atmosphere


13

Fig. 2.5. Schematic diagrams of various operations of direct SC process: (a) pouring melt into

die cavity, (b) solidification under pressure and (c) removal of casting [10].

Fig. 2.6. Schematic diagram of cold chamber HPDC [10].


14

is not maintained in the cold chamber HPDC process and the molten Mg is poured into a shot

cylinder either manually, or automatically. The plunger then injects it fast (5–10 m/s) into the

die cavity, where solidification takes place under high pressure (35–140 MPa). Finally, the

casting is removed. The entire sequence takes typically around 1 min. The cold chamber HPDC

process is shown in Fig. 2.6. High production rate is ensured in this process [10].

2.6 Tensile and creep behaviour of Mg alloys in different casting conditions

Gutman et al. [11] studied the influence of casting parameters on creep and strength of the

AZ91D alloy fabricated by permanent mould cast as well as die-cast, and concluded better

creep resistance in permanent mould cast alloy by β-Mg17Al12 phase precipitation. In die-cast

alloy, the creep resistance was controlled by the presence of micro and macro porosity,

whereas, the β-Mg17Al12 phase precipitation determined the creep resistance of permanent

mould alloy. Variation in tensile properties of the AZ91 alloy synthesized by sand-casting with

varying solidification rate was reported by Caceres et al. [12]. The solidification rate was varied

with the help of cast iron chills placed at one end of the cast plate and riser at the other end.

The grain size and dendrite arm spacing of the alloy decreased towards the chill end on account

of higher solidification rate. However, some aspects of microstructure and mechanical

properties relationship were inadequately understood by them like effect of intermetallic

particles on ductility, Hall-Petch effect and thickness effects in HPDC alloys. Dargusch et al.

[13] investigated the effect of varying section thicknesses of a HPDC AZ91 alloy and

concluded that the mechanical properties at elevated temperature improved as section thickness

was decreased. This was attributed to the fine grain size on the skin region of the thinner

castings. Han et al. [14] reported the increase in creep resistance of the AC52 alloy with

increased cooling rate. Slower cooling rates resulted in coarsening of the eutectic phases.

However, they observed best creep resistance in the alloy with intermediate cooling rate as a


15

result of combined effect of dispersion and solid solution strengthening. Zhu et al. [15] too

studied the effect of varying cooling rate using chill at one end of the cast plate on the fraction

of β-Mg17Al12 phase in the sand cast AZ91 alloy and they observed that the fraction of eutectic

β-Mg17Al12 decreased and discontinuous β-Mg17Al12 increased towards the chill end of the

casting. However, they did not investigate the mechanical properties of the alloy. Superior

mechanical properties including creep behaviour were reported in the die-cast Mg-3Sm-0.5Zn-

0.4Zr (wt.%) alloy by Zheng et al. [16], whereas, ingot-cast alloy showed better tensile

properties at 200°C. Better creep resistance of the die-cast alloy was attributed to the dynamic

precipitation of fine platelets in the matrix of α-Mg, which resulted in hindering the dislocation

motion during creep deformation. Ferri et al. [17] also noticed improvement in strength of

ZAXLa05413 alloy subjected to unidirectional cooling as compared to the isothermal cooling.

The higher rate of heat extraction resulting in microstructural refinement attributed to enhanced

mechanical properties. Kim et al. [18] too asserted that the Mg-4Al-2Sn alloy in die-cast

condition exhibited better creep resistance than that in IC condition. The grain refinement and

finer intermetallic phases lead to superior creep resistance in the die-cast alloy. Bai et al. [19]

also observed the difference in tensile and creep properties with varying casting condition in

the Mg-4Al-(1-4) La alloy. The Mg17Al12 phase was supressed on increase of La addition

resulting in better creep resistance. Mondal et al. [20] reported better creep resistance in the

HPDC MRI230D alloy as compared to that of the IC alloy. The better creep resistance in the

HPDC alloy in their investigation was attributed to the presence of denser and finer network of

the C36 phase leading to effective dispersion strengthening.


16

2.7 Corrosion behaviour of Mg alloys in different casting conditions

Song et al. [21] investigated influence of microstructure on corrosion performance of the die-

cast AZ91D alloy and reported better corrosion resistance in the die-cast alloy as compared to

the high purity crucible cooled AZ91 alloy. Ambat et al. [22] too concluded higher corrosion

resistance and better passivation behaviour in the die-cast AZ91D alloy than in the IC alloy on

account of fine grain structure and β-Mg17Al12 phase. Mathieu et al. [23] investigated corrosion

behaviour of the AZ91 alloy and concluded that Al content of the α-Mg solid solution played

a major role in corrosion resistance. Ballerini et al. [24] in their work on die-cast and sand cast

AZ91D alloy claimed that the finer microstructure resulted in better corrosion resistance of the

die-cast alloy. However, reverse trend was also observed by Blawert et al. [25] in their

investigation on SC and thixocast MgZnRE alloys, i.e., thixocast alloy with coarser grain size

exhibited better corrosion resistance. Casting defects like porosity affects corrosion

performance of Mg alloys considerably. The effect of porosity on corrosion performance of

Mg alloys was investigated by Song et al. [21,26]. They reported that the effective surface area

may perhaps be considerably greater than the surface area exposed and due to confined

exchange of solution between inside and outside of the pores, corrosion may increase

considerably inside the pores. Cao et al. [27] too investigated effect of porosity on the corrosion

behaviour of the Mg-0.1Si alloy and deduced that corrosion rate increased due to the presence

of porosities.

17

Chapter 3. Experimental procedure

The AZ91 alloy is fabricated by three different casting techniques. Several characterization

techniques have been employed to investigate the microstructures and properties of these

alloys. In this chapter, the details of fabrication processes and characterization of these alloys

are described.

3.1 Fabrication of the alloys

The IC alloy was produced by melting commercially available AZ91D Mg alloy ingots with

nominal chemical composition of Mg-9Al-0.7Zn-0.2Mn (wt.%) in an electrical resistance

furnace inside a stainless steel crucible of capacity 1 Kg. Melting of the alloy was carried out

at 720°C up to 90 min. A mixture of SF6 (0.5 vol.%) and Ar (99.5 vol.%) was purged during

the whole process to protect the melt from oxidation. To ensure uniform distribution of the

alloying constituents, the melt was stirred for 10 min with a holding time of 3 min to allow the

oxide particles and other inclusions to settle at the bottom of the crucible. The melt was poured

at a pouring temperature of 700°C into a steel mould (dimension: 50 mm diameter and 250 mm

height) preheated to 210°C and allowed to cool till it attained room temperature. The same

melting practice was employed to fabricate the SC alloy. Solidification took place in around

60 s under a pressure of 120 MPa applied directly on the melt by a vertical hydraulic press till

the completion of solidification. The actual photograph of the setup employed in the present

investigation is shown in Fig. 3.1. The HPDC alloy was cast in a cold chamber HPDC machine

(Model: FRECH DAK 450-54RC) in the form of cylindrical rod of 19 mm diameter and 179

mm length. The melt temperature and the pressure during solidification was maintained at

690°C and 300 bar, respectively.


18

Fig. 3.1. Actual photograph of the SC setup employed to fabricate the IC and SC alloys.


19

3.2 Specimen preparation

Standard metallographic procedure was followed to prepare the specimens for microstructural

analysis. The specimens were etched by acetic picral (a solution of 100 ml ethanol, 20 ml

distilled water, 10 ml acetic acid and 6 gm picric acid). Tensile specimens proportional to

standard dimensions were prepared in accordance with the ASTM E8M-03 standard. For creep

testing, specimens of dimension of 6 mm diameter and 15 mm length were machined from the

IC, SC and HPDC alloys by wire electrical discharge machining (EDM). Specimens for

electrochemical investigation were polished successively on 600, 800, 1000, 2000 and 3000

grit SiC emery paper followed by cleaning in distilled water and ultrasonic cleaning in ethanol.

3.3 Characterization of the alloys

3.3.1 X-ray diffraction (XRD)

The phases present in the alloys were identified by an X-ray diffractometer (XRD) (Model:

Rigaku Japan/Ultima-IV) employing CuKα (λ = 1.541 Å) radiation. A scan rate of 2º/min in

the scan range of 10-90º with step size of 0.05 was employed.

3.3.2 Optical microscopy (OM)

Microstructures of the alloys were observed using an optical microscopy (Model: ZEISS

AxioVision I10) equipped with an AxioVision image analyser software. The ASTM E1382

standard was followed for grain size measurement of the alloys in three different casting

conditions. The grain size, volume fraction of phases, secondary dendrite arm spacing (SDAS)

and porosity were quantified with the help of image analyser.


20

3.3.3 Scanning electron microscopy (SEM)

Microstructures of the alloys were observed using scanning electron microscopy (SEM)

(Model: FEI Quanta 200) equipped with an energy dispersive X-ray spectroscopy (Model:

Oxford Instruments). Microstructures of tensile and creep tested specimens were also observed

under SEM in order to understand the underlying deformation mechanisms.

3.3.4 Mechanical characterization

3.3.4.1 Tensile test

Tensile specimens proportional to standard dimensions were prepared and tested in accordance

with the ASTM E8M-03 standard in a universal testing machine (Model: INSTRON 5967) at

ambient (≈ 25°C), 150 and 200°C temperature.

3.3.4.2 Creep test

Creep tests were carried out using a lever arm (10:1) creep setup (Model: ATS 2330) in

compression at a stress of 70 MPa and temperature of 150ºC for 50 h.

3.3.5 Corrosion test

Electrochemical corrosion tests were performed using Gill AC potentiostat/galvanostat

equipped with sequencer software. The setup was a conventional electrochemical cell equipped

with three electrodes. It consisted of the AZ91 alloy specimen as working electrode, Pt counter

electrode and Ag/AgCl (3 mol/L KCl) reference electrode. The tests were carried out at ambient

temperature in 0.5 wt.% NaCl solution at neutral pH. One experiment consisted of three tests

as listed below.


21

i. OCP of each specimen was recorded for a duration of 30 min.

ii. After recording OCP, electrochemical impedance spectroscopy (EIS) was performed

by the application of a sinusoidal signal of amplitude 5 mV (root-mean-square value)

in the frequency range of 100 kHz to 0.01 Hz. The measurements were repeated after

every 2 h.

iii. Potentiodynamic polarization scans were carried out at -200 mV with respect to the free

corrosion potential with a scanning rate of 0.2 mV/s. The tests were terminated when

the current density reached 1 mA/cm2.

After completion of the corrosion tests, the specimens were cleaned by distilled water for

observation under microscope. The loose and non-adherent corrosion products were collected

for further analysis by XRD and Fourier Transform Infrared Spectroscopy (FTIR) (Model:

Shimadzu IRPrestige-21). FTIR spectra were recorded at ambient temperature in the mid-IR

region of 4000–400 cm-1. The adherent corrosion products from the specimens were removed

by using a solution prepared according to ASTM G1-03 standard (20 g chromic acid, 2 g

barium nitrate and 1 g silver nitrate dissolved in distilled water to yield 100 ml solution). The

specimens were immersed in the solution for about 60 s. They were immediately washed in

distilled water and ultrasonically cleaned in ethanol, and dried. The specimens were cut

perpendicular to the corroded surface in order to measure the depth affected by corrosion.

22

Chapter 4. Microstructural characterization

In this chapter the detailed microstructural characteristics of all the fabricated AZ91 alloys have

been discussed.

4.1 As-cast microstructure

Fig. 4.1 shows the XRD patterns obtained from the AZ91 alloy in three different casting

conditions. It is evident from the figure that the alloy in all three casting conditions consists of

primary Mg (α-Mg) and β-Mg17Al12 phases. Intensity of the peak corresponding to β-Mg17Al12

phase was the highest and lowest in the IC and HPDC alloys, respectively. Thus, the volume

fraction of the low melting point β-Mg17Al12 phase was the highest and lowest in the IC and

HPDC alloys, respectively. The presence of relatively lower amount of β-Mg17Al12 phase in

the HPDC alloy is expected to improve its creep resistance.

Optical micrographs of the AZ91 alloy in all three casting conditions are shown in Fig. 4.2.

The microstructures of all the alloys consist of α-Mg phase and bright β-Mg17Al12 phase present

along grain boundaries and triple points. The alloy in all casting conditions consists of

polygonal grains. The average grain size determined by linear intercept method was 70.26 ±

4.53, 41.20 ± 3.85 and 26.44 ± 1.96 μm for the IC, SC and HPDC alloys, respectively. Thus,

the HPDC alloy exhibited relatively finer grain size as compared to that of the IC and SC alloys.

Slow cooling rate associated with IC allowed ample time for grain growth, whereas, in the

other two alloys the cooling rate was sufficiently higher resulting in relatively finer grain size.

The high degree of undercooling associated with the faster cooling rate following HPDC


23

20 30 40 50 60 70 80 90

(c)

(b)

-Mg

-Mg17

Al12

Re

lative

In

ten

sity (

A. U

.)

2(Degree)

(a) IC

(b) SC

(c) HPDC

(a)

Fig. 4.1. XRD patterns obtained from the AZ91 alloy in three different casting conditions.


24

Fig. 4.2. Optical micrographs of the (a) IC, (b) SC and (c) HPDC alloys.


25

facilitated the formation of a large number of nuclei in the liquid Mg alloy. However, the

sluggish diffusion rate at the corresponding temperature leading to low growth rate resulted in

microstructural refinement [28]. The SDAS was measured by linear intercept method and the

values were 68.01 ± 6.53, 35.50 ± 5.09 and 31.99 ± 4.67 μm for the IC, SC and HPDC alloys.

It is obvious that the SDAS decreased with increase in solidification rate.

SEM micrographs of the AZ91 alloy in all three casting conditions are shown in Fig. 4.3. EDS

analysis carried out in the grain interior of the IC alloy shown in Fig. 4.3(a) exhibited an

average composition of 94.86 ± 0.61 Mg, 4.72 ± 0.44 Al, 0.29 ± 0.11 Zn and 0.13 ± 0.05 Mn

(at.%), which corresponds to α-Mg. Similarly, EDS analysis of the bright lamellar phase

present along grain boundaries and triple points revealed an average composition of 61.44 ±

5.6 Mg and 41.74 ± 4.49 Al (at.%). Therefore, the grain boundary phase is β-Mg17Al12 phase

as marked by the arrows in Fig. 4.3. In addition, a few globular Al-Mn rich particles (shown

by arrows in Fig. 4.3(b)) were also observed and EDS analysis exhibited an average

composition of 44.02 ± 4.6 Al, 24.28 ± 2.1 Mn and 31.41 ± 1.4 Mg (at.%). EDS analysis was

also carried out on the other two alloys revealing the grain interior as α-Mg; the bright phase

present along grain boundaries and triple points as β-Mg17Al12 phase; and a few Al-Mn-rich

phase. Thus, the existence of same phases i.e., α-Mg, β-Mg17Al12 and Al-Mn-rich phases in the

AZ91 alloy in all three casting conditions was confirmed.

The IC, SC and HPDC alloys differed much on the following grounds. The calculated average

volume fraction of the grain boundary phase in the IC alloy (i.e., 17.88 ± 0.24%) was higher as

compared to that observed in the SC (i.e., 8.16 ± 3.15%) and HPDC (i.e., 7.59 ± 0.45%) alloys.

Thus, the fraction of β-Mg17Al12 phase decreased with increase in solidification rate owing to


26

Fig. 4.3. SEM micrographs of the (a) IC, (b) SC and (c) HPDC alloys.


27

the fact that high cooling rate did not allow Al in α-Mg matrix to precipitate and therefore, it

formed Mg-Al-Zn solid solution during solidification. Barbagallo et al. [29] too observed the

decreased volume fraction of the grain boundary β-Mg17Al12 phase with increase in cooling

rate in the AZ91 alloy. Another noteworthy dissimilarity observed was the difference in

morphology of the eutectic β-Mg17Al12 phase, which is quite sensitive to small variations in

cooling conditions and alloy content. The observed eutectic phase was partially divorced,

coarse and discontinuous in the IC alloy. On the other hand, it was partially divorced and

constituted a continuous network along the grain boundaries and the triple points in the SC

alloy; fully divorced and discontinuous in the HPDC alloy. In addition, the IC alloy exhibited

wider eutectic structure as compared to that of the SC and HPDC alloys. The former alloy

allowed sufficient time for growth of the precipitates, thereby widening the structure of the

eutectic phase. Nave et al. [30] too reported similar microstructural characteristics i.e., the β-

Mg17Al12 phase was partially divorced in the IC Mg-9Al alloy. The reason behind the formation

of divorced eutectic in Mg alloys is well explained in the literature [30,31]. A binary Mg-Al

alloy of eutectic composition (i.e., Mg-33Al (wt.%)) typically solidifies with a regular lamellar

or fibrous morphology comprising of α-Mg and β-Mg17Al12 phases depending on cooling rates.

The tendency to form divorced eutectic is increased with increase in volume fraction of the

primary phase resulting from low Al content in the alloy. Thus, hypoeutectic Mg alloy with

low solute concentration generally exhibits divorced or partially divorced eutectic

microstructures [31]. The tendency to form divorced eutectic was enhanced by addition of

small amount of Zn that strongly segregated into the liquid as compared to Al. Further, this

tendency increased with the increase in cooling rate. In all alloys the α-Mg phase appeared as

light grey and it gradually became darker towards grain boundaries due to higher concentration

of Al and Zn that was detected by EDS analyses. A similar observation was made in the IC

AZ91 alloy by Nave et al. too [30].


28

In order to assess the segregation tendency of Al, the concentration of Al interior of α-Mg near

grain boundary was measured using EDS and the values were 6.29 ± 1.11, 9.39 ± 0.36 and

14.62 ± 5.13 (at.%) in IC, SC and HPDC alloys. Thus, the Al concentration interior of α-Mg

in the alloy in three different casting conditions was not same and it was the highest in the

HPDC alloy and the lowest in the IC alloy, as expected. The excess solubility of Al within α-

Mg grain is expected to exhibit difference in lattice parameter of the AZ91 alloy in three

different casting conditions. The rapid cooling rate associated with the HPDC process was

responsible for the formation of supersaturated solid solution of Al in α-Mg grains.

Consequently, the amount of Al available for the β-Mg17Al12 formation was lower and

therefore, the final volume fraction of the β-Mg17Al12 phase in the HPDC alloy at the end of

solidification was relatively less.

Table 4.1. Quantitative analysis of the microstructural features of all the alloys.

Alloy IC SC HPDC

Grain size (μm) 70.26 ± 4.53 41.20 ± 3.85 26.44 ± 1.96

Secondary dendrite arm

spacing (SDAS) (μm)

68.01 ± 6.53 35.50 ± 5.09 31.99 ± 4.67

Volume fraction of

β-Mg17Al12 (%)

17.88 ± 0.24 8.16 ± 3.15 7.59 ± 0.45

Al content in α-Mg

(at.%)

6.29 ± 1.11 9.39 ± 0.36 14.62 ± 5.13

Density (gm/cm3) 1.816 1.916 1.568


29

Another notable distinction witnessed was the difference in porosity content. The density

values of the IC, SC and HPDC alloys were 1.816, 1.916 and 1.568 gm/cm3, respectively. Thus,

the porosity (marked by arrows in Fig. 4.2(c)) content was the highest in the HPDC alloy and

it was not significant in the SC alloy with the IC alloy exhibiting the intermediate porosity. The

higher porosity observed in the HPDC alloy was attributed to the difference in liquid metal

temperature (i.e., casting temperature), die (mold) temperature, metal injection rate into the die

as well as casting conditions including the process of melt preparation and the possibility of air

entrainment during metal injection into the die [32]. In addition, the distinct difference in the

solidification pressure among the three casting technologies employed in the present

investigation might also influenced the microstructures of the three alloys.

30

Chapter 5. Mechanical characterization

In this chapter the tensile and creep properties of the AZ91 alloy fabricated by three different

casting techniques are compared and correlated with their microstructures. Observation of

microstructures of the tested specimens is also carried out to understand the underlying

mechanisms of deformation.

5.1 Tensile behaviour

Typical stress vs. strain curves of the SC alloy tested at ambient, 150 and 200°C temperature

is shown in Fig. 5.1 as a representative plot for all the alloys. Fig. 5.2(a) shows the tensile

properties derived from the stress vs. strain curves of the AZ91 alloy in three different casting

conditions tested at ambient temperature. It is observed that the casting condition significantly

influenced the tensile properties of the AZ91 alloy. The 0.2% proof stress (YS), ultimate tensile

strength (UTS) and ductility (% Elongation) were the lowest for the IC alloy, intermediate for

the HPDC alloy and the highest for the SC alloy.

The YS and UTS values decreased and correspondingly the ductility increased for the alloy in

all casting conditions when tested at 150 and 200°C temperature, as shown in Fig. 5.2(b&c).

The ductility of the IC, SC and HPDC alloys increased by 21.0, 14.7 and 15.7% at 150°C; and

by 26.3, 25.0 and 20.6% at 200°C, respectively from the ambient temperature. The best tensile

properties were observed in the SC alloy at all temperature levels employed in the present

investigation. It exhibited 3.4, 3.9 and 4.1% ductility with corresponding YS of 140, 132 and

120 MPa at ambient, 150 and 200°C temperature. The improved tensile behaviour of the SC

alloy could be associated with relatively finer grain size and negligible porosity present in the

casting.


31

0.00 0.01 0.02 0.03 0.04 0.050

20

40

60

80

100

120

140

160

3

2

Eng

g. str

ess (

MP

a)

Engg. strain (mm/mm)

1. Ambient temperature

2. 150°C

3. 200°C

1

Fig. 5.1. Stress vs. strain plots of the SC alloy tested at ambient, 150°C and 200°C

temperature.


32

78

89 9298

103

140

1.9

1.6

3.4

IC HPDC SC0

25

50

75

100

125

150

Str

en

gth

(M

Pa

)

Alloy

YS

UTS

Elongation

(a)

0

1

2

3

4

5

Elo

ng

atio

n (

%)

7178 80

9488

132

2.3

1.9

3.9

IC HPDC SC0

25

50

75

100

125

150

Str

en

gth

(M

Pa

)

Alloy

YS

UTS

Elongation

(b)

0

1

2

3

4

5

Elo

ng

atio

n (

%)

6673 76

8690

120

2.4

2.0

4.1

IC HPDC SC0

25

50

75

100

125

150

Str

en

gth

(M

Pa

)

Alloy

YS

UTS

Elongation

0

1

2

3

4

5

Elo

ng

atio

n (

%)

(c)

Fig. 5.2. Comparison of tensile properties of the IC, SC and HPDC alloys at (a) ambient, (b)

150°C and (c) 200°C temperature.


33

The relationship between YS (𝜎0.2) and average grain size (d) is expressed by Hall-Petch

equation as follows:

𝜎0.2 = 𝜎0 + kd−1/2 (1)

where 𝜎0 is the friction stress and k is the locking parameter, and they both depend on the

material [9]. The YS of all the alloys were calculated theoretically based on the grain size

employing Hall-Petch relation and the values are presented in Table 5.1. For calculation, k =

142.2 MPa√µm and 𝜎0 = 73 MPa were considered for the AZ91 alloy [33]. It is evident that

the calculated values of the YS were slightly higher than that determined by experiment. The

discrepancy in YS values might be attributed to several reasons. Firstly, the values of the

k and σ0 were corresponding to HPDC AZ91 alloy and not corresponding to IC and SC alloys.

Secondly, the microporosities present in the HPDC alloy might contributed negatively to

experimental YS of the alloys. The calculated YS of the SC alloy was in good agreement with

the experimentally obtained YS owing to the presence of negligible porosity in it.

Table 5.1. Comparison of the calculated and experimentally determined YS of all the alloys.

Alloy Calculated

grain size (µm)

Calculated

YS (MPa)

Experimentally

determined YS (MPa)

IC 70.26 ± 4.53 90 78

SC 41.20 ± 3.85 95 92

HPDC 26.44 ± 1.96 101 89

Evidently, the IC alloy revealed relatively low YS, UTS and ductility due to the presence of

higher volume fraction of brittle β-Mg17Al12 phase rendering the alloy brittle. The coarse grain

size also contributed to the inferior tensile properties of the IC alloy. The relatively poor tensile

properties of the HPDC alloy might be due to structural inhomogeneity comprising of


34

microporosities present in the casting. Dúl et al. [34] too in their investigation observed

inhomogeneity in the internal structure of the HPDC alloy. The existing pores where lattice of

the alloy was significantly distorted acted as stress raisers and crack initiated. Once the crack

formed, it connected the neighbouring pores and thus, allowing easy path for rapid crack

propagation.

Fig. 5.2(a-c) shows the micrographs of the fracture surfaces of the tensile specimens of the IC,

SC and HPDC alloys tested at ambient temperature. Features revealed are cracks (shown by

A), porosity (B), quasi-cleavage plane (C), steps (D), plastic deformation zone (E), shallow

dimples (F) and these were marked in the micrographs. The micrograph of the IC alloy

exhibited the presence of various cleavage steps. Slight deformation zone were also present.

Multiple cracks were observed in the micrograph of the HPDC alloy and these cracks were

initiated from the sharp edges of the pores. A few quasi-cleavage planes (C) with steps of

various sizes were also observed. These are the characteristic of brittle intergranular cleavage

fracture that took place along the grain boundaries. Owing to the limited number of slip systems

operated at ambient temperature, dislocation movement in the neighbouring grains was

restricted leading to reduced ductility of all the alloys. Quasi-cleavage fracture was the failure

mode in the SC alloy in which the cracks initiated and propagated locally forming a shallow

dimple like morphology on the fracture surface. Presence of second phases and their nature

played an important role behind the origin of micro cracks. During tensile test micro cracks

might initiated at the interfaces of brittle eutectic present along the grain boundaries, thus

rendering the alloy brittle. Patel et al. [35] also observed secondary cracks on the fracture

surface propagated along the interface of eutectic phase in the AZ91D alloy. The presence of

relatively more number of shallow dimples and plastic deformation zones on the fracture

surface of the SC alloy evidenced its higher ductility. Fracture surfaces of all the specimens of


35

the alloys tested at 150 and 200°C exhibited quasi-cleavage fracture as well. Mg alloy generally

exhibits limited ductility owing to limited slip systems operating at ambient temperature.

However beyond 225°C ductility of Mg alloy enhances owing to activation of additional slip

systems containing prismatic and pyramidal planes along which cross-slip takes place [36].

The elevated testing temperatures (i.e., 150 and 200°C) employed in the present investigation

were below 225°C and accordingly, additional slip systems were not completely activated,

which resulted less improvement in ductility.


36

Fig. 5.3. Fracture surfaces of the tensile specimens of the (a) IC, (b) SC and (c) HPDC alloys

tested at ambient temperature.


37

5.2 Creep behaviour

The characteristic creep curves (strain vs. time) of the AZ91 alloy in IC, SC and HPDC

conditions tested at 70 MPa stress and temperature of 150°C for a duration of 50 h are shown

in Fig. 5.4(a). The strain rate vs. time curves derived from the strain vs. time curves in Fig.

5.4(a) is shown in Fig. 5.4(b). Apparently, different cooling rates associated with the IC, SC

and HPDC alloys resulted varied creep characteristics. All the curves displayed a distinct

primary (transient) creep followed by a secondary (steady-state) creep. All the specimens were

tested under compression and therefore, the tertiary stage was not observed in the plots. The

creep rate (strain rate) (𝜀̇) was calculated from the steady-state part of the strain vs. time curves

for all the alloys by fitting straight line as shown in Fig. 5.4(c). The IC, SC and HPDC alloys

exhibited a creep rate of 7.26×10-8, 1.60×10-8 and 9.56×10-8 s−1, respectively. Obviously, the

creep rate was the highest in the HPDC, intermediate in the IC and the lowest in the SC alloys

in the present investigation. The creep rate of the HPDC and IC alloys; IC and SC alloys; HPDC

and SC alloys varies by a factor of 1.32, 5.97 and 4.53, respectively. Thus, the casting routes

significantly affected creep behaviour of the AZ91 alloy.

The observed differences in creep behaviour amongst the IC, SC and HPDC alloys were

attributed to several reasons. Different cooling conditions prevailing during castings resulted

varied microstructural features leading to the variation in creep behaviour. The difference in

amount and morphology of the β-Mg17Al12 phase present along grain boundaries of the IC, SC

and HPDC alloys resulted in the variation. It is speculated based on the review of literature on

creep behaviour of Mg alloys and especially of AZ91 alloy that dislocation creep is the


38

0 40000 80000 120000 160000 2000000.000

0.004

0.008

0.012

0.016

0.020

0.024

Str

ain

Time (s)

1. HPDC

2. IC

3. SC

1

2

3

(a)

0 40000 80000 120000 160000 20000010

-9

10-8

10-7

10-6

10-5

(b)

1. HPDC

2. IC

3. SC

Str

ain

ra

te (

s-1)

Time (s)

1

2

3

0 40000 80000 120000 160000 2000000.000

0.004

0.008

0.012

0.016

0.020

0.024

(c)

1.6010-87.2610-8

3. y = 1.60E-08x + 3.05E-03

R² = 9.87E-01

2. y = 7.26E-08x + 3.86E-03

R² = 9.99E-01

Str

ain

Time (s)

1. HPDC

2. IC

3. SC

1. y = 9.56E-08x + 3.07E-03

R² = 9.99E-01

9.5610-8

έ (s-1)

Fig. 5.4. (a) Strain vs. time plots for the IC, SC and HPDC alloys creep tested at 70 MPa

stress and temperature of 150°C for 50 h; (b) Strain rate vs. time plots and (c) Calculation of

strain rates by linear fit to the plots shown in (a).


39

proposed creep mechanism in the present alloy in all three casting conditions at the stress and

temperature level employed [20,37,38]. Creep resistance increases if the dislocation glide

and/or dislocation climb faces hindrance on the basal and/or non-basal plane leading to

increased work hardening. The presence of discontinuous β-Mg17Al12 phase in the HPDC alloy

could not effectively provide obstacle to dislocation motion. The SEM micrograph of the

HPDC alloy creep tested at a stress of 70 MPa and temperature of 150ºC is shown in Fig. 5.4.

It is evident that the existing porosities in the alloy joined together providing easy path for

crack initiation and growth during creep test. A similar observation was also made by Mondal

et al. [20] on the creep tested die-cast MRI230D alloy. Therefore, higher porosity content of

the HPDC alloy deteriorated its creep resistance. Thus, the extent of porosity played a crucial

role in the creep behaviour of the HPDC alloy, which was verified by Gutman et al. [11,32] as

well. They observed that the minimum creep rate under the comparable condition increased

considerably with increase in porosity. As a consequence, easy crack initiation and growth

during the creep test took place, reducing creep performance of the HPDC alloy. On the

contrary, presence of negligible porosity and continuous network of eutectic phase in the SC

alloy contributed to its superior creep resistance. Amberger et al. [39] too concluded that the

continuous and interconnected network of the intermetallic phase enhanced creep resistance in

the modified AZ91 and MRI230D alloys. The intermediate creep performance exhibited by the

IC alloy was attributed to the presence of relatively lesser porosity as well as coarser and

discontinuous eutectic phase. The creep resistance of the IC alloy was better than that of the

HPDC alloy owing to the relatively lower porosity content. On the other hand, it was slightly

inferior as compared to that of the SC alloy due to the presence of coarser and discontinuous

eutectic phase. The coarse and discontinuous eutectic phase contributed negatively to creep

resistance due to lower pinning of dislocations. Kim et al. [18] also concluded that coarse

eutectic phase favoured the formation of cracks allowing easy dislocation movement.


40

Fig. 5.5. SEM micrograph of the HPDC alloy creep tested at 70 MPa stress and temperature

of 150°C for 50 h.

Porosity

41

Chapter 6. Corrosion behaviour

This chapter presents the electrochemical corrosion behaviour of all the alloys. The observed

corrosion behaviour is correlated with their microstructures.

6.1 Corrosion response

6.1.1 Open circuit potential (OCP)

Fig. 6.1 shows the OCP of all the three alloys recorded for 30 min. OCP of the SC alloy initially

increased continuously with increase in time i.e., it shifted towards more noble potential, which

indicates that the passive film is continuously growing with time. It then drops off indicating

sudden breakdown of the film and continued with fluctuations implying several cycles of

formation and subsequent dissolution of the passive film. OCP of the IC alloy too initially

increased slowly and after that it became almost stable. On the other hand, OCP of the HPDC

alloy initially decreased followed by a slight increase with increase in time. The fluctuations in

OCP values of these two alloys were relatively less as compared to that of the SC alloy. Thus,

the passive films formed on the surfaces of the IC and HPDC alloys were relatively more

protective.

Average values of OCP were calculated and these were -1476, -1494 and -1494 mV for the

HPDC, SC and IC alloys, respectively as shown in Fig. 6.2. The characteristic corrosion

behavior of the alloys cannot be predicted from OCP values, however, it serves as a reference

for the subsequent tests i.e., EIS and potentiodynamic polarization scan.


42

0 400 800 1200 1600 2000-1600

-1550

-1500

-1450

-1400

3

2

Pote

ntia

l (m

V v

s. A

g/A

gC

l)

Time (Sec)

1. HPDC

2. SC

3. IC

1

Fig. 6.1. Variation of OCP with time of the HPDC, SC and IC alloys.

HPDC SC IC-1600

-1550

-1500

-1450

-1400

OC

P (

mV

vs. A

g/A

gC

l)

Alloy

Fig. 6.2. Comparison of OCP of the HPDC, SC and IC alloys.


43

6.1.2 Electrochemical impedance spectroscopy (EIS)

Nyquist plots of all the three alloys are shown in Fig. 6.3. For comparison the data

corresponding to 0 and 10 h for all the alloys are plotted separately and shown in Fig. 6.4. The

HPDC and IC alloys exhibited the best and worst long-term corrosion resistance with the SC

alloy exhibiting intermediate corrosion rate. The Nyquist plots in all the cases consists of two

capacitive loops and one inductive loop. The high-frequency (HF) capacitive loop indicates

charge transfer resistance (Rct) from the surface to the electrolyte. In other words, the HF loop

can be considered as the diffusion of ions through passive film consisting of oxides or

hydroxides of Mg. The low-frequency (LF) capacitive and inductive loops indicated localized

corrosion attack on the surface. Considering a simple Randles circuit model for the long-term

corrosion behaviour of the alloys, the highest and lowest polarization resistance were exhibited

by the HPDC and IC alloys, respectively. Thus, passive film formed on the surface of the IC

alloy was not protective due to localized corrosion attack. Higher polarization resistance in the

HPDC and SC alloys resulted in greater deposition of corrosion products on the β-Mg17Al12

phase and lesser dissolution of α-Mg phase. This indicated weak galvanic coupling in the SC

and HPDC alloys. Diameter of capacitive loops obtained for all the alloys progressively

decreased, which implies that corrosion resistance decreased with increase in time i.e., from 0

to 10 h, as shown in Fig. 6.3 and 6.4. The same behaviour as observed in Nyquist plots was

also observed in Bode magnitude plots (BMPs) and Bode phase plots (BPPs) shown in Fig. 6.5

and 6.6, respectively. Positive slope at low-frequency (LF) depicted inductive behaviour,

whereas, negative slope indicated capacitive behaviour in BMP. Reverse is true in BPPs i.e.,

negative slope at LF represented inductive behaviour and positive slope at HF depicted

capacitive behaviour. Straight lines parallel to frequency-axis in BMPs and BPPs indicated

resistive behaviour. Drastic decrease in impedance was observed at the LF region of the BMPs


44

for all the alloys, as shown in Fig. 6.5. The breakpoint frequencies for all the sequences for all

the alloys were nearly same indicating constant rate of hydrogen emission [40].


45

0 500 1000 1500 2000 2500-400

-200

0

200

400

600

800

(a)

Z"

(oh

m.c

m²)

Z' (ohm.cm²)

0 h 2 h 4 h 6 h 8 h 10 h

0 500 1000 1500 2000 2500 3000 3500-400

0

400

800

1200

1600

(b)

Z"

(oh

m.c

m²)

Z' (ohm.cm²)

0 h 2 h 4 h 6 h 8 h 10 h

0 100 200 300 400 500-25

0

25

50

75

100

125

150

(c)

Z"

(oh

m.c

m²)

Z' (ohm.cm²)

0 h 2 h 4 h 6 h 8 h 10 h

Fig. 6.3. Nyquist plots of (a) HPDC, (b) SC and (c) IC alloys.


46

0 500 1000 1500 2000 2500 3000 3500-400

0

400

800

1200

1600

Z"

(oh

m.c

m²)

Z' (ohm.cm²)

HPDC

SC

IC

(a)

0 500 1000 1500 2000 2500 3000 3500-400

0

400

800

1200

1600

(b)

Z"

(oh

m.c

m²)

Z' (ohm.cm²)

HPDC

SC

IC

0 200 400 600 800 1000 1200 1400 1600-260

-130

0

130

260

390

520

650

(c)

Z"

(oh

m.c

m²)

Z' (ohm.cm²)

HPDC

SC

IC

Fig. 6.4. Nyquist plots of all the alloys at (a) 0 h, (b) 10 h and (c) magnified view of (b).


47

10-2

10-1

100

101

102

103

104

105

102

103

(a)

Imp

ed

an

ce

(o

hm

.cm

²)

Frequency (Hz)

0 h 2 h 4 h 6 h 8 h 10 h

10-2

10-1

100

101

102

103

104

105

102

103

(b)

Imp

ed

an

ce

(o

hm

.cm

²)

Frequency (Hz)

0 h 2 h 4 h 6 h 8 h 10 h

10-2

10-1

100

101

102

103

104

105

102

103

(c) 0 h 2 h 4 h 6 h 8 h 10 h

Imp

ed

an

ce

(o

hm

.cm

²)

Frequency (Hz)

Fig. 6.5. Bode impedance plots of the (a) HPDC, (b) SC and (c) IC alloys.


48

10-2

10-1

100

101

102

103

104

105

-10

0

10

20

30

40

50

Ph

ase

an

gle

()

Frequency (Hz)

0 h 2 h 4 h 6 h 8 h 10 h

(a)

10-2

10-1

100

101

102

103

104

105

-10

0

10

20

30

40

50

60

70

(b)

Ph

ase

an

gle

()

Frequency (Hz)

0 h 2 h 4 h 6 h 8 h 10 h

10-2

10-1

100

101

102

103

104

105

-10

0

10

20

30

40

50

(c)

Ph

ase

an

gle

()

Frequency (Hz)

0 h 2 h 4 h 6 h 8 h 10 h

Fig. 6.6. Bode phase angle plots of the (a) HPDC, (b) SC and (c) IC alloys.


49

6.1.3 Potentiodynamic polarization scan

The results of potentiodynamic polarization scans for all the alloys are shown in Fig. 6.7.

Corrosion rates were determined by Tafel extrapolation method from the cathodic branch of

the polarization curves and shown in Fig. 6.8. Corrosion rates were 0.52, 2.97 and 11.71 mm/y

for the HPDC, SC and IC alloys, respectively. The values of various parameters obtained from

the plots shown in Fig. 6.7 are summarized in Table 6.1. The HPDC alloy exhibited more

positive Ecorr and the lowest corrosion rate owing to higher Al content in the α-Mg phase.

However, among the three alloys the IC alloy containing lowest concentration of Al in the α-

Mg phase revealed intermediate corrosion resistance. The maximum shift in Ecorr from OCP

towards more positive potential indicated the best corrosion resistance of the HPDC alloy

among the alloys employed in the present investigation.

Table 6.1. Values of various parameters obtained from the potentiodynamic polarization

plots.

Alloy Ecorr

(mV) a

(mV) c

(mV) Icorr

(mA/cm2) Corrosion rate

(mm/y)

HPDC -1306 58 392 0.02 0.52

SC -1382 171 380 0.13 2.97

IC -1420 262 386 0.51 11.71


50

-1700 -1600 -1500 -1400 -1300 -1200

10-4

10-3

10-2

10-1

100

101

1

2

Cu

rre

nt d

en

sity (

mA

/cm

2)

Potential (mV vs. Ag/AgCl)

1. HPDC

2. SC

3. IC3

Fig. 6.7. Potentiodynamic polarization plots for all the alloys.

HPDC SC IC0

2

4

6

8

10

12

14

Co

rro

sio

n r

ate

(m

m/y

)

Alloy

Fig. 6.8. Comparison of corrosion rate obtained from Tafel extrapolation.


51

6.2 Microstructural observation after corrosion tests

Representative micrographs of passive films formed on the corroded surfaces of all the alloys

are shown in Fig. 6.9. It is obvious from Fig. 6.9(a) that the surface of the HPDC alloy was

least affected by corrosion. The surface of the IC alloy shown in Fig. 6.9(c) exhibited

discontinuous passive film with large number of cracks and microvoids on it. In contrast, the

film on the SC alloy was relatively uniform and almost free from microvoids as shown in Fig.

6.9(b).

The XRD pattern obtained from the corroded surface of the SC alloy is shown in Fig. 6.10 as

a representative picture for all the alloys. The pattern consists of Mg(OH)2 phase as corrosion

product along with the α-Mg and β-Mg17Al12 phases. The non-adherent corrosion product

obtained from the corroded specimen was collected for FTIR analysis. The representative FTIR

pattern obtained from the corrosion product of the SC alloy is shown in Fig. 6.11. Presence of

Mg(OH)2 (brucite) and Mg5(CO3)4(OH)2.4(H2O) (hydromagnesite) was confirmed in the

corrosion product [41]. The very distinct –OH peaks at 3698 and 473 cm−1 confirmed the

presence of Mg(OH)2. The sharp CO32− asymmetric bending peaks at 1484 and 1426 cm−1 was

attributed to hydromagnesite. The peak at 854 cm−1 was due to the CO32− bending vibration

band. The peak at 3451 cm−1 was also observed in the reference spectra of hydromagnesite.


52

Fig 6.9. SEM micrographs of the surfaces corresponding to the (a) HPDC, (b) SC and (c) IC

alloys before removal of corrosion products.


53

10 20 30 40 50 60 70 80 90

In

ten

sity (

A. U

.)

2 (Degree)

Mg

Mg17

Al12

Mg(OH)2

20 40 60 80

Mg

Mg17

Al12

Inte

nsity (

A. U

.)2 (Degree)

Fig. 6.10. XRD pattern obtained from the corroded surface of the SC alloy.

4000 3500 3000 2500 2000 1500 1000 5000.0

0.5

1.0

1.5

2.0

2.5

3.0

Ab

so

rba

nce

(A

. U

.)

Wavenumber (cm-1)

-OH bond

-CO2-

3 bond

Fig. 6.11. FTIR pattern of corrosion product obtained from the corroded surface of the SC

alloy.


54

Fig. 6.12(a-c) shows the corroded surfaces of all the alloys after removal of corrosion products.

It is evident that the α-Mg phase was preferentially dissolved and a network of β-Mg17Al12

phase is standing as a corrosion resistant barrier in all the alloys. SEM micrograph of cross-

section of the corroded surface of the IC alloy is shown in Fig. 6.13. It is obvious that the β-

Mg17Al12 phase remained as a corrosion resistant barrier and the α-Mg phase was selectively

removed by corrosion. Thus, preferential dissolution of the α-Mg phase was also confirmed

from the observation of cross-section. Corrosion initiated at the interface of α-Mg and β-

Mg17Al12 phases, which is reported in literature as well [45]. There is a potential difference

between the α-Mg and β-Mg17Al12 phases resulting in micro-galvanic coupling between them

[42-44]. The β-Mg17Al12 phase is cathodic to α-Mg. Thus, the α-Mg phase was removed by

galvanic corrosion. The β-Mg17Al12 phase acted as a barrier to corrosion up to some extent,

however, it was undermined owing to the selective dissolution of α-Mg phase. Thus, the

superior corrosion resistance exhibited by the HPDC alloy in the present investigation was

attributed to: the higher concentration of Al in the α-Mg phase; the observed smallest grain size

in it. The beneficial effects of finer grain size to corrosion resistance of Mg alloys was well

explained elsewhere [21,46].


55

Fig. 6.12. SEM micrographs of the (a) HPDC, (b) SC and (c) IC alloys after removal of

corrosion products.


56

Fig. 6.13. SEM micrograph of cross-section of the corroded surface of the IC alloy.

57

Chapter 7. Conclusions

The relationships of microstructure with tensile, creep and corrosion behaviour of the AZ91

Mg alloy produced by three different casting techniques was investigated. The following

conclusions are drawn:

i. All the as-cast alloys consisted of primary Mg (α-Mg) and β-Mg17Al12 phases. The

volume fraction of the β-Mg17Al12 phase was the highest in the ingot casting (IC),

intermediate in the squeeze-casting (SC) and the lowest in the high pressure die-casting

(HPDC).

ii. The best tensile properties were exhibited by the SC alloy at all the temperatures

employed in the present investigation owing to the presence of negligible porosity and

relatively finer grain size. Fracture surfaces of the broken tensile specimens revealed

quasi-cleavage fracture.

iii. The best and worst creep resistance were exhibited by the SC and HPDC alloys with

the IC alloy exhibiting the intermediate creep resistance. Poor creep resistance of the

HPDC alloy was attributed to the presence of higher amount of porosity that allowed

easy crack initiation and growth during creep test. The presence of negligible porosity

and continuous network of eutectic phase contributed to superior creep resistance of the

SC alloy.

iv. The best and worst corrosion resistance were exhibited by the HPDC and IC alloys.

The superior corrosion resistance of the HPDC alloy was attributed to the relatively

finer grain size and higher percentage of Al content in the α-Mg phase.

Chapter 7. Conclusions

58

v. The SC alloy exhibited the best tensile and creep properties with intermediate corrosion

resistance among the alloys employed in the present investigation. Therefore, the SC

technique is recommended for mass production of the AZ91 alloy.

59

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62

Bio-Data

Name Hrishikesh Shastri

Date of birth 25 August 1990

Permanent home address Plot No. 1200/13, Kelo Vihar Colony,

Raigarh, Chhattisgarh-496001.

Email [email protected]

Educational qualification

Course Institution Board /University Year of

Completion

Performance

% / CGPA

B.E.

(Metallurgy)

O.P. Jindal Institute of

Technology, Raigarh

(C.G.)

Chhattisgarh Swami

Vivekanand Technical

University, Bhilai, (C.G).

2012 8.53

12th O.P. Jindal School,

Raigarh CBSE 2008 73.3

10th O.P. Jindal School,

Raigarh CBSE 2006 85.6

Proposed publications from this work

1. H. Shastri, A. K. Mondal, K. Dutta, H. Dieringa and S. Kumar, Correlation of

microstructure with tensile and creep behaviour of AZ91 Mg alloy fabricated by three

different casting techniques (Manuscript under preparation)

2. H. Shastri, A. K. Mondal, K. Dutta, H. Dieringa and S. Kumar, Correlation of

microstructure with corrosion behaviour of AZ91 Mg alloy fabricated by three different

casting techniques (Manuscript under preparation)

Correlation of microstructure with tensile, creep and corrosion ...

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