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Editorial Panel 

Paolo ANTONA 

Materials Consultant. 

Pietro APPENDINO 

Professor, Material Technology and Applied Chemistry; Faculty of Engineering, 

Turin Polytechnic. 

Marcello BADIALI 

R&D Director, Teksid Aluminum. 

Giuseppe CAGLIOTI 

Professor, Solid State Physics ; Faculty of Engineering, Milan Polytechnic. 

Enrico EVANGELISTA 

Professor, Metallurgy; Faculty of Engineering, Ancona Polytechnic. 

Luca Paolo FERRONI 

Marketing Director, Teksid Aluminum. 

Merton C. FLEMINGS 

Head Department of Materials Science and Engineering, M.I.T., Cambridge, Mass. 

Sergio GALLO 

Invited Professor, Applied Chemistry & Metallurgy; Faculty of Engineering, Turin 

Polytechnic. Past President Teksid S.p.A. 

Cinzia MODENA 

Editorial Coordinator, Metallurgical Science & Technology, Teksid Aluminum. 

Claudio MUS 

Materials Consultant. 

Walter NICODEMI 

Professor, Siderurgy & Siderurgical Technology; Faculty of Engineering, Milan Polytechnic. 

Editor in chief: Marcello BADIALI 

Registrazione presso il Tribunale di Torino n. 3298 del 12 maggio 1983 

Associato alla Unione Stampa Periodica Italiana 

Authors are asked to supply two typewritten copies of their papers. These 

should be laid out in accordance with the “Instructions for authors” on the 

inside back cover. All correspondence should be addressed to: 

Segreteria di redazione 

Metallurgical Science and Technology 

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e-mail: journal@teksidaluminum.com 

Publication distributed free of charge 

© Copyright 1983 Teksid S.p.A. - All rights reserved. 

Editorial coordination: Cinzia Modena, Marketing, Teksid Aluminum 

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Printed in Italy: Graficat - Via Cuniberti, 47 - Torino 

No part of the texts published in this journal may be reproduced, whether in 

the original language or in translation, without permission in writing from 

Teksid Aluminum. 

Papers submitted for publication as a rule should illustrate unpublished original 

research works, experiments, or critical reviews. Publication will depend on 

the approval of the Editorial Panel.

E. Gariboldi, D. Ripamonti, 

L. Signorelli, G. Vimercati, 

F. Casaro 

S. Seifeddine, T. Sjögren, 

I. L. Svensson 

M. El Mehtedi, L. Balloni, 

S. Spigarelli, E. Evangelista, 

G. Rosen, B.H. Lee, C.S. Lee 

Metallurgical 

Science and Technology 

A journal published 

by Teksid Aluminum 

twice a year 

Vol. 25 No. 1, July 2007 

FRACTURE TOUGHNESS AND 

MICROSTRUCTURE IN AA 2XXX 

ALUMINIUM ALLOYS 

VARIATIONS IN MICROSTRUCTURE AND 

MECHANICAL PROPERTIES OF CAST 

ALUMINIUM EN AC 43100 ALLOY 

COMPARATIVE STUDY OF HIGH 

TEMPERATURE WORKABILITY OF ZM21 

AND AZ31 MAGNESIUM ALLOYS 

PAGE 

3 

12 

23

EDITORIAL 

The present issue marks an important goal for Teksid Aluminum: the 25th 

anniversary of Metallurgical Science and Technology (MS&T). Twenty five years 

of life represent a remarkable milestone which is not commonly reached 

by a scientific journal. MS&T is one of the few specialized editorial products 

edited by an industrial company. It has been always delivered world wide, 

free of charge and with continuity. 

This moment deserves a word on the nature of the Journal, on its intrinsic, 

permanent features as well as on its evolution and future perspectives. 

The birth of MS&T originated from the initiative of professor Sergio Gallo, 

former President of Teksid S.p.A., together with full professors Aurelio 

Burdese of the Politecnico di Torino and Walter Nicodemi of the Politecnico 

di Milano. They suggested to earmark part of the financial resources destined 

to the image promotion activities of the company for a specialized scientific 

review. Their proposal was warmly agreed by Antonio Mosconi, the then 

Managing Director of Teksid, who, in the first issue’s introduction, wrote: 

“The aim of Metallurgical Science and Technology is to initiate a dialogue, 

already part of the scene in other countries, between all those 

resources that are committed to the search for new frontiers in 

metallurgical techniques, so as to derive the greatest synergies from 

them. 

Metallurgical Science and Technology then enters the scene as an 

instrument whereby the dialogue between metallurgical science and 

metalworking industry can be taken to greater depths. 

‘Academic’ metallurgy indeed, has often shown signs of hankering 

after pure science, estranged from industry. It should not be forgotten 

that its final objective remains practical application. 

A journal such as this is not as ambitious as to imaging that it can 

entirely fill the goal. Its aim is rather to inculcate the habit of 

comparison and so cast a permanent bridge between the two parties. 

Its very existence shows that the need to set up a more systematic 

and continuous two-way flow of information is seen by both sides as 

an indispensable prelude to its development.” 

Since these targets and perspectives were indicated, several changes have 

occurred in the life of Teksid Group, the sponsor and editor of the journal. 

The original name of the company, Teksid, evokes a core business centered 

on cast iron foundry. Founded in 1978 as a spin-off of Fiat metallurgical 

activities, its experience rooted in the very beginning of the Italian industrial 

development which started-up in the early century. Steel and cast-iron 

production were then the company’s core business focus. The global trend 

toward “energy conservation” triggered by the recurrent oil crises of the 

seventies and eighties induced Teksid to share and favor the efforts aiming 

at vehicle weight reduction: the interest of the Company then gradually 


shifted to light metal alloys – like magnesium and, 

above all, aluminum – as base materials. Teksid was 

then split into divisions depending on the material 

their industrial activity was dedicated to (iron, 

aluminum and magnesium). In 2002 the aluminum 

division was sold to the private equity market to 

give birth to Teksid Aluminum, today’s publisher 

of MS&T. 

It is known that today’s market dynamics, especially 

in the car production industry, is much and 

dramatically faster than in the past years, and Teksid 

Aluminum had to react to the changing scenario 

through a severe restructuring and divestiture 

process, which largely resized its traditional 

footprint of an international and industry leading 

company. 

The evolution experienced by Teksid Aluminum 

and the radical changes occurred during the last 

decades in science, in technology and consequently 

in our society, have left the objectives of the journal 

unchanged, and its contents have evolved 

coherently with the scientific and technologic 

trends of metallurgy. 

In the Company strategy, these changes have been 

accompanied by a constant search for innovative 

and competitive product solutions, by a constant 

effort to enable the costumers enjoy the company’ 

s collaboration, from the concept stage to the 

delivery of unfinished or completely machined 

products and by a constant attention the best 

combination of technical performance and 

product quality. 

This evolution has been reflected by MS&T too. 

This is not the place for presenting an exhaustive 

list of the topics covered by the journal in all 

these years. However probably some readers 

will remember academic and technological 

contributions dedicated to innovative forming 

techniques – such as lost-foam or semisolid 

forming, – or to the challenges derived by the 

emergence of new structural materials such as cast

magnesium alloys, or to thermomechanical 

treatments – like liquid hot isostatic pressing, – 

and to new experimental methods for testing metal 

properties and performance – thermoelastoplastic 

yield stress vs. conventional yield stress, and much 

more. 

Since its birth, this journal has hosted contributions 

of both academic research and metalworking 

industry from all over the world. Perhaps the 

‘academicians’ have cooperated more 

enthusiastically than the industrial researchers to 

the “two-way flow of information” originally 

wished, but these attitudes is to be attributed both 

to the confidential nature of industrial research 

and to the old academic caveat “publish or perish” 

still alive in the web era.. 

In all these years Teksid Aluminum has treasured the MS&T bridging role 

between science and industrial expertise in the R&D programs of the Group. 

Furthermore the metallurgic community could enjoy a 25-year long support 

to scientific divulgation. We can assert that the targets, originally indicated 

by Mr. Mosconi, have been reached. Nowadays MS&T papers published since 

1983 can be found in university libraries, academic institutions and industrial 

R&D laboratories all over the world. Notably, all the issues published since 

the year 2000 can be also retrieved comfortably from the Company website 

(www.teksidaluminum.com). 

The topics such as those mentioned above and unpredictable new ones – 

specifically related to light metals and alloys - will continue to arouse interest 

of contributors and readers of MS&T for many years to come, as the 

increasing cooperation and subscriptions testify. 

In a joyful anniversary like this one, we cannot conclude this editorial other 

than by congratulating with Metallurgical Science & Technology for its longevity 

and scientific maturity, and by wishing Happy Birthday and a long life! 


Editorial Panel

FRACTURE TOUGHNESS AND 

MICROSTRUCTURE IN AA 2XXX 

ALUMINIUM ALLOYS 

1 E. Gariboldi, 1 D. Ripamonti, 1 L. Signorelli, 1 G. Vimercati, 2 F. Casaro 

1 Politecnico di Milano, Dipartimento di Meccanica, Milano (MI) 

2 Varian Vacuum Technologies, Leinì (TO) 

Abstract 

The paper presents the toughness properties of forgings made of two AA 2xxx series 

aluminium alloys with different microstructural conditions. Fracture toughness tests in 

crack opening mode I were performed on compact tension specimens machined from 

the forgings in different orientations. The tests were performed both at room temperature 

and at 130°C. 

Fracture toughness properties were related to microstructural and fractographic features 

of the alloys in order to discuss on their failure mechanisms. The effect of the coarse 

intermetallic phases within grains or at their boundaries in the different conditions was 

underlined. The testing temperature, within the range here investigated, neither affected 

fracture toughness properties nor failure mechanisms. 

KEYWORDS 

Al alloy AA2014, Al alloy AA2618, fracture toughness, microstructure. 

3 - Metallurgical Science and Technology 

Riassunto 

Il lavoro illustra le proprietà di tenacità alla frattura misurate 

in forgiati di grandi dimensioni realizzati con due leghe di 

alluminio della serie 2xxx in differenti condizioni 

microstrutturali. Dai forgiati sono stati ricavati provini CT 

con differenti orientazioni sui quali sono state condotte 

prove di tenacità a frattura secondo il modo I di apertura 

della cricca. Le proprietà ottenute sono state correlate 

con la microstruttura riscontrata nei campioni e completate 

con analisi frattografiche atte ad individuare i meccanismi 

di cedimento. È stato così messo in luce l’effetto delle 

diverse microstrutture ed in particolare delle particelle 

grossolane di fasi intermetalliche presenti a bordo grano 

o all’interno dei grani che differenziano i forgiati nelle leghe 

esaminate di composizione più complessa.

1. INTRODUCTION 

Aluminium-alloy forging is currently used to 

manufacture structural components of relatively 

large and complex shape. The plastic deformation 

imparted to the material can positively affect its 

microstructure by promoting recrystallization 

cycles and a greater homogeneity of alloying 

elements. However, it should be considered that 

in large size forgings, the relatively low amount of 

plastic strain given to the alloy cannot completely 

refine the structure and intermetallic particles as 

in other small-size wrought products such as 

extruded bars or rolled sheets. In addition, the 

slower quenching rates experienced by large 

forgings result in lower mechanical properties 

achieved after the subsequent aging process. Large 

differences in cooling rate between surface and 

centre of large forgings during solution annealing 

also result in remarkable residual stresses, that are 

often relieved by inserting a plastic deformation 

step after quenching and before the aging 

treatment [1]. This method also modifies the 

precipitation sequence and kinetics of the alloy. 

The above described effects significantly affect the 

tensile and fracture properties of aluminium alloy 

forgings. Focusing the attention on the fracture of 

aluminium alloys, it was reported that toughness 

is strictly related to the presence of coarse 

particles, 0.1 to 10 µm in diameter, that can be 

either non-equilibrium particles formed during 

2. MATERIALS INVESTIGATED 

The present investigation was carried out on three 

forgings having a roughly cylindrical shape with a 

diameter of 250 mm, made of aluminium 

alloys AA2014 (Al4CuSiMg) and AA2618 

(Al2Cu1.5MgNi). The parts had been forged from 

extruded bars of diameter 190 mm with different 

manufacturing cycles. 

Two forged samples of the AA2014 alloy were 

produced by forging in two steps at 390°C. The 

samples were then heat treated to T6 temper by 

different parameters. A first sample, hereafter 

referred to as 2014-A forging was solution 

annealed at 505°C for 6 hours, water quenched 

and artificially aged at 160°C for 14 hours, following 

the usual industrial heat treatment route. In the 

case of the forging in AA2618 alloy (hereafter 

referred to as 2618 forging), the solution annealing 

at the usual temperature for this material, 530°C, 

lasted 1 hour, and it was followed by water 

quenching and by artificial aging at 190°C for 20 

solidification or inclusions from insoluble impurities [2]. These particles 

crack easily as the matrix deforms within the plastic flow zone at the crack 

tip and causes the typical ductile fracture mode where crack propagates 

via coalescence of voids. The amount, size and distribution of these second 

phase particles is thus relevant for the fracture toughness properties and 

even material with comparable tensile properties can display significantly 

different fracture toughness properties. 

In addition to the abovementioned effect, the role played by submicrometer 

particles (0.01 to 0.5 µm in diameter) need to be considered. In the case of 

the same volume fraction and microstructural features of coarse particles, 

a substantial modification of toughness can be observed in age hardenable 

aluminium alloys varying the amount and characteristics of fine hardening 

particles. The behaviour is complex and the fine hardening particles are 

responsible for it. It is well known that the presence of particle having 

suitable distribution and size enhances the resistance of peak aged alloys to 

deformation and thus tends to reduce the extension of the plastic zone, 

positively affecting toughness [2, 3]. On the other hand, the lower strainhardening 

capacity of the material in the peak aged condition with respect 

to the underaged condition, also gives rise to local plastic instabilities that 

significantly contribute to reduce the material toughness [2]. Further, where 

grain boundary precipitate free zones are observed, strain localization in 

these regions and intergranular ductile fracture can occur [3, 4]. In these 

cases the fracture toughness depends on the spacing and size of the voidnucleating 

particles at grain boundaries [4, 5]. The higher fracture toughness 

displayed by alloys in the underaged with respect to over-aged condition as 

well as the transition towards intergranular ductile fracture mode as 

overaging proceeds confirm this latter effect [2, 4]. 

The aim of the present paper is to contribute to a better understanding of 

the correlation between microstructure, tensile and toughness properties 

of aluminium forgings as a result of different thermomechanical cycles, 

focusing in particular to the role of large second phase particles. 

hours, following a common industrial practice. The 2014-B forging was 

produced and heat treated as for 2618 material, leading to a substantially 

overaged matrix and to intermetallic phases distribution rather different 

than that of 2014-A forging. 

Light optical microscopy observations and Vickers hardness tests (0.98N 

load) were performed to evaluate the general microstructural features. 

Tensile test specimens were machined from the forgings in the hoop direction 

in regions characterized by a homogeneous structure and hardness. 

Two sets of Compact Tension (CT) fracture toughness specimens were 

machined with cracks laying in diametral planes of the forgings. In the first 

set, the crack propagation direction was radial (CR direction according to 

ASTM E399-90 and B645-02 [6, 7]) while in the second set it was longitudinal 

(CL direction). The specimens had a thickness B of 20 mm [6, 7]. 

Tests were carried out on a MTS 810 universal testing machine, equipped 

with an environmental chamber suitable for test temperatures up to 250°C. 

Before fracture toughness tests, a precracking stage was performed at test 

temperature until a total crack length (machined notch and fatigue crack) 

of about 20 mm was reached. Precracking was performed under load control 

(sinusoidal cycles at 10 Hz frequency) while crack length was monitored 

via the elastic compliance technique by measuring the Crack Opening 

Displacement (COD). The stress intensity factor during pre-cracking 

decreased linearly with crack length from 11 to 8 MPa•m 1/2 . Fracture 

4 - Metallurgical Science and Technology

toughness tests were carried out imposing a displacement rate of 0,025 

mm/s, monitoring the COD and applied load until the occurrence of unstable 

crack propagation. 

In a second stage of the research study, the toughness of the materials at 

130°C was investigated. The J IC integral was measured according to the 

single-specimen technique and the ASTM E1820-01 standard [8]. The J IC 

parameter was evaluated instead of K IC since the material toughness was 

expected to be greater than that measured at room temperature. The tests 

were performed monitoring crack length via the compliance method at 

fixed steps of COD. In order to reduce the strain accumulated under load 

during each load step due to creep effects at 130°C, the holding period 

under constant applied load (P) for the adjustment of the crack length was 

fixed to 1 s. J Q was estimated according to the mentioned standard as the 

intercept between the power-law curve (fitting the experimental data within 

the range stated by ASTM E1820-01) and the straight line passing to the 

3. RESULTS 

3.1 MICROSTRUCTURE 

All forgings were characterized by grains elongated along the main plastic 

flow path experienced during forging. Coarse intermetallic particles were 

also present in the microstructure, as expected for these alloys. More 

specifically, the 2014-A material (figure 1a) was characterized by elongated 

grains with a transversal size of about 50 µm and by large intermetallic 

particles aligned in the flow direction: globular Al 2 Cu (θ) particles (bright 

particles in figure 1a) and blocky shaped clustered particles containing Fe, 

Mn, Si and Cu (darker particles in the same figure). In regions where these 

particles were observed small equiaxed grains were often detected 

A B 

C D 


point ∆a = 0.2 mm, P = 0 N with a slope 

corresponding to a double flow stress (this latter 

corresponding to the average between the yield 

and the ultimate tensile stress). 

Validity requirements of J 1C tests were not met in 

all cases since unstable crack propagation or popin 

phenomena occurred in some samples. In these 

cases the K Q parameter was evaluated from the 

P-COD curves. 

Fracture surfaces were observed by scanning 

electron microscopy (SEM). Metallographic 

sections were also cut perpendicularly to the crack 

plane to observe the crack propagation path by 

optical microscopy and SEM. 

suggesting the local recrystallization effect induced 

by these intermetallics. 

The direction of plastic flow was less evident in 

the 2014-B samples, that displayed rather elongated 

grains of 50-100 µm in transverse size together 

with smaller equiaxed grains in the regions 

containing large intermetallic particles (figure 1b). 

These were in lower amount with respect to the 

previous alloy. Particles with dark appearance at 

grain boundaries (figure 1b) were of complex 

chemical composition. On the contrary globular 

Al 2 Cu particles had a bright appearance in 

Fig. 1: 

Typical microstructure of 

the investigated forgings. 

2014-A (a), 2014-B (b), 

2618 at low (c) and 

higher (d) magnification. 

The longitudinal axis of 

forgings is horizontal in 

all the micrographs.

micrographs and were observed both inter- and 

intragranularly. 

The 2618 sample was characterized by very large 

grains slightly elongated in the flow direction. Their 

size in the transverse direction was of the order 

of 500 µm, as shown in figure 1c. As expected, 

considering the composition of the AA2618 alloy 

[1], several intragranular FeNiAl 9 particles were 

observed together with globular Al 2 Cu and (CuFe) 

Al 3 precipitates. 

The microstructure and hardness of the forgings 

were in all cases homogeneous in the regions 

where tension and toughness specimens were 

sampled. 

3.2 MECHANICAL BEHAVIOUR 

The average tensile properties in hoop direction 

of the investigated materials at room temperature 

and at 130°C are summarized in Table I. At room 

temperature, the 2014-A forging showed the 

highest tensile strength and the lowest ductility. 

The increase in test temperature from 20 to 130°C 

led to a significant reduction in the ultimate tensile 

strength (UTS) and 0,2% proof stress (YS) and to 

a corresponding significant increase in reduction 

of area (Z) and elongation (A) at fracture. 

The results of the fracture toughness tests are 

summarized in Table 2 while representative load 

vs. COD graphs are shown in Figure 2. Here, the 

lines needed for the evaluation of K Q are added 

A) B) 


Table 1: Tensile properties of 

the investigated alloys 

YS [MPa] UTS [MPa] A [%] Z [%] 

Temperature 20°C 130°C 20°C 130°C 20°C 130°C 20°C 130°C 

2014-A 421 386 454 402 1,8 9,3 3,7 15,8 

2014-B 357 341 425 393 7,3 8,4 14,2 16,9 

2618 350 345 416 409 6,4 8,6 10,8 12,3 

Table 2: Fracture toughness of the materials 

investigated (K Q or K IC according to ASTM B645 

and K JIC accordiing to ASTM E1820) 

Specimen Temperature 2014-A 2014-B 2618 

direction [°C] MPa . m Ω MPa . m Ω MPa . m Ω 

CL 20 19,3 (K IC ) 23,0 (K Q ) 22,4 (K Q ) 

CL 20 20,1 (K IC ) 24,9 (K Q ) 23,3 (K Q ) 

Average CL 20 19,7 (K IC ) 24,0 (K Q ) 22,9 (K Q ) 

CL 130 18,8 (K IC ) 23,4 (K JIC ) 21,9 (K Q ) 

CR 20 23,4 (K IC ) 24,8 (K Q ) 26,9 (K Q ) 

CR 20 22,8 (K IC ) 24,4 (K Q ) 26,7 (K Q ) 

Average CR 20 23,1 (K IC ) 24,6 (K Q ) 26,8 (K Q ) 

CR 130 21,4 (K JIC ) 23,4 (K JIC ) 25,6 (K Q ) 

to the experimental curves 

The condition of plain strain crack propagation in CT specimen fracture 

can be met for sufficiently large specimens. The minimum specimen thickness 

Fig. 2: Load vs. COD plots of the samples tested at 20°C. 

A) 2014-A forging, CL specimen; B) 2014-B forging, CR specimen

Table 3: F values for the materials and specimen 

orientation investigated 


direction [°C] 

CL 20 9.4 4.8 4.9 

CL 20 8.7 4.1 4.5 

Average CL 20 9.1 4.5 4.7 

CL 130 9.1 4.2 4.9 

CR 20 6.4 4.1 3.4 

CR 20 6.8 4.3 3.4 

Average CR 20 6.6 4.2 3.4 

CR 130 7.4 3.3 3.6 

Table 4: J IC values (N/mm) of the investigated 

forgings tested at 130°C 


direction [°C] [N/mm] [N/mm] [N/mm] 

CL 130 - 7.23 - 

CR 130 5.96 7.00 - 

(B) to be adopted can be calculated, at the end of the tests, according to 

the two standard above considered: 

B min =5·(K Q /YS) 2 according to ASTM B645 (1) 

B min =2,5·(K Q /YS) 2 according to ASTM E399 (2) 

A) B) 

Fig. 2: Load vs. COD plots of the samples tested at 20°C. 

A) 2014-A forging, CL specimen; B) 2014-B forging, CR specimen 


Thus, the minimum thickness required by the E399 

standard is half of that required by the B645 

standard for aluminium alloys. The agreement to 

plain strain conditions and the possible deviation 

from the minimum value of B can be evaluated 

using a parameter F, defined as: 

F=B·(YS/K Q ) 2 (3) 

Where F greater than 5 means that the 

requirements for plain strain condition are met 

for both standards and valid K IC can be obtained. 

2,5

according to the ASTM E1820 standard. As far as 

the AA2618 alloy that resulted to be the toughest 

material at room temperature, at 130°C it showed 

unstable crack propagation (Figure 3b) that 

prevented the evaluation of J IC. In the cases of valid 

J IC , K JIC was computed, assuming plain strain 

conditions, by using the following correlation 

proposed by ASTM E1820 standard: 

3.3. FRACTOGRAPHIC 

OBSERVATIONS 

Fractographic observations were performed at 

two main locations on fracture surface of each 

specimen in the crack propagation region: the first 

next to, the latter at 4 mm from the boundary 

with fatigue precrack. Selected images at low 

magnification of the first location for the different 

alloys and specimen orientation are reported in 

figure 4. In the unstable crack propagation region, 

the 2014-A forging showed the typical features of 

ductile fracture, generated by nucleation of dimples 

mainly from the coarser intermetallic particles 

fractured in a brittle way (figure 5a). Within the 

matrix, regions with dimples of far smaller size 

were also visible, laying on well-defined planes 

A) 

C) 

K JIC = [(E/(1-v 2 ))•J IC ] 0.5 (4) 

K JIC values are also listed in Table 2 for comparison purpose. Examination of 

this table clearly shows that toughness of 2014-B forging is greater than 

that of the same alloy in forging 2014-A. Further, it can be stated that in all 

the examined samples, crack propagation in CL orientation is favoured 

with respect to that in CR orientation. 

(figure 5b). These can be correlated to the presence of fine particles that in 

some cases were observed inside the dimples in high magnification images 

that suggest ductile intergranular fracture. 

The fracture surfaces of the 2014-B forging did not show significant 

differences from 2014-A material, with relatively extended regions of 

microdimples (figure 7a) that could be correlated to the presence of coarse 

grain boundary particles (figure 6b). The size and distribution of these 

particles corresponded to those observed at grain boundaries (figure 4c). 

The 2618 forging, showed a transgranular ductile fracture mode (figures 

4d, 7b). Dimples observed on fracture surfaces nucleated from the 

homogeneously distributed FeNiAl 9 particles (figures 6c and 6d). 

The fracture surfaces of specimens tested at 130°C of forgings made of 

AA2014 alloy were similar to those tested at room temperatures (figure 

8a). On the contrary, as the temperature increased, small-size microdimples 

were observed on the fracture surface of the forging made of AA2618 alloy 

(compare figure 7b and 8b). 

Fig. 4: Fracture surface of cracks propagated at room temperature in the boundary region between fatigue precracking and unstable crack 

propagation region. 2014-A forging, CL (a) and CR (b) directions. 2014-B forging; CR direction (c) and 2618, CR direction (d). 

B) 

D) 


A) B) 

A) 

C) 

Fig. 5: Fracture surface of cracks propagating at room temperature in 2014-A forging in CL (a) and CR (b) direction. 

Fig. 6: Microstructure in the region of unstable crack propagation on longitudinal section of CT specimens. a) 2014-A forging, CL direction; 

b) 2014-B forging, CL direction; c) and d) 2618 forging, CL and CR directions, respectively. 

B) 

D) 

A) B) 

Fig. 7: Fracture surface sampled in CL direction of 2014-B forging (a) and of 2618 forging (b) tested at room temperature. 


A) B) 

Fig. 8: Fracture surface of cracks propagated at 130°C in specimen with CL orientation of 2014-A (a) and 2618 (b) forgings. 

4. DISCUSSION 

In the examined forgings crack propagated in a 

ductile manner, in some cases by an intergranular 

ductile mode, linking the early fractured coarse 

intermetallic particles, located at inter or 

intragranular position in the different materials 

investigated. 

Following the approach proposed by Hahn and 

Rosenfield [2] for aluminium alloys, it can be stated 

that crack propagates when the size of extensive 

plastic strain formed ahead of the crack tip 

corresponds to the average coarse interparticle 

distance (δ) Further, according to these authors, 

K and d are correlated as follows: 

δ =(0.5K 2 )/(E•YS) (5) 

In the case of 2618 forging the δ values estimated 

using room temperature tensile and toughness 

properties were 11 and 15 µm for specimen 

orientations CL and CR, respectively, comparable 

to the average interparticle size between the 

intragranular particles (mainly FeNiAl 9 particles) 

in longitudinal and radial direction, respectively. The 

preferred particle orientation and their tendency 

to align axially in the forging regions where 

specimens were sampled corresponds to a greater 

CONCLUSIONS 

The toughness properties in opening mode I of 

forgings made of AA2014 and AA2618 aluminium 

alloys with different microstrctural conditions were 

presented. 

Fracture toughness ranged between 19 and 26 

MPam 0.5 , depending on the material and specimen 

direction. Fracture toughness was in general lower 

interparticle distance in the crack propagation direction of CR specimens. 

As previously presented, in the specimens obtained from forgings made of 

AA2014 alloy the crack proceeds in a ductile manner but along an 

intergranular path, intercepting the oriented coarse aligned intermetallic 

particles axially. The values of δ computed from equation 5 for these two 

forgings are 6.7 and 9 µm for CL and CR specimen of 2014-A forging, 

respectively. In the case of the last forging (2014-B), δ equals 12 µm for 

both CL and CR directions. Thus, also in the case of AA2014 alloy, d is 

comparable to the interparticle distance along the crack path. It can be 

thus stated that fracture toughness is correlated to the mean interparticle 

distance along the crack path. 

The above observations well agree with the simple correlation proposed 

by Hahn and Rosenfield [2] for the case of forged Al-alloy parts (where 

fracture is initiated by cracking of large intermetallic particles once reached 

critical stress/strain levels in the region of extensive deformation at the 

crack tip [3]) and fracture toughness is correlated to the mean interparticle 

distance along the crack path. 

Thus, forgings of the same heat treatable alloy, even in the same heat 

treatment condition and with comparable hardness and tensile properties 

can show significantly different fracture behaviour, depending on intermetallic 

particle population. Especially, as the mean interparticle distance along the 

crack path between the coarser intermetallic particles increases, the 

nucleation of voids at these particles is shifted at higher applied loads and 

toughness is improved. 

The role of coarse intermetallic particles and the need to reduce their 

amount, to optimize their size and distribution in forgings (taking into account 

the most critical crack propagation directions in the components) in order 

to increase material toughness is therefore highlighted. 

for specimen sampled in CL direction. In the examined forgings crack 

propagated in ductile manner, in some cases by an intergranular ductile 

mode, linking in any case the voids formed at the early fractured inter- or 

intragranular coarse intermetallic particles. No significant difference in 

toughness nor in the fracture mode was observed when tests were 

performed at room temperature or at 130°C. 

Fracture toughness properties were related to the presence and distribution 

of the coarse intermetallic phases within grains or at their boundaries in 

the different investigated alloys. 


The simple correlation proposed by Hahn and Rosenfield between fracture 

toughness and the mean interparticle size was found to be applicable when 

the interparticle distance along the crack path (different for different sampling 

direction in forgings where these particle were aligned along flow directions) 

is taken into account. An observation and quantitative assessment of the 

REFERENCES 

[1] J.S. Robinson, R.L. Cudd, J.T. Evans. Creep resistant aluminium alloys 

and their applications. Material Science and Technology, 19 (2003) pp. 

143-155. 

[2] G.T. Hahn, A.R. Rosenfield. Metallurgical Factors Affecting Fracture 

Toughness of Aluminum Alloys. Metall. Trans. A, (1975) pp. 653-668. 

[3] T. Kobayashi. Strength and fracture of aluminium alloys. Materials Science 

Engineering, A286 (2000) pp. 333-341. 

[4] A.P. Reynolds, E.P. Crooks. The effect of thermal exposure on the fracture 

behaviour of aluminium alloys intended for elevated temperature 

service. in “Elevated temperature on fatigue and fracture”, ASTM STP 

1297. Williamsbourg, Virginia(USA), (1997) 191-205. 

[5] B.Q. Li, A.P. Reynolds. Correlation of grain-boundaries precipitates 

parameters with fracture youghness in an Al-Cu-Mg-Ag alloy subjected 

to long-term thermal exposure. Journal of Materials Science, 33 (1998) 

pp. 5849-5853. 

[6] ASTM B 645-02 Practice for plane-strain fracture toughness testing of 

aluminium alloys. 

[7] ASTM E 399-99 Standard test method for plane-strain fracture 

toughness of metallic materials. 

[8] ASTM E1820-01. Standard test method for measurement of fracture 

toughness. 


distribution of coarse intermetallic phases on 

optical micrographs can be a useful tool to check 

fracture toughness properties of different forgings 

in their peak aged condition when yield strength 

is known.

VARIATIONS IN MICROSTRUCTURE AND 

MECHANICAL PROPERTIES OF CAST 

ALUMINIUM EN AC 43100 ALLOY 

Salem Seifeddine, Torsten Sjögren and Ingvar L Svensson 

Jönköping University, School of Engineering, Component technology - Sweden 

Abstract 

The microstructure and mechanical properties of a gravity 

die and sand cast Al-10%Si-0.4%Mg alloy, which is one of 

the most important and frequently used industrial casting 

alloys, were examined. Tensile test samples were prepared 

from fan blades and sectioned through three positions 

which experienced different cooling rates. Furthermore, 

the inherent strength potential of the alloy was revealed 

by producing homogeneous and well fed specimens with a 

variety of microstructural coarseness, low content of oxide 

films and micro-porosity defects, solidified in a laboratory 

environment by gradient solidification technology. The 

solidification behaviour of the alloy was characterized by 

thermal analysis. By means of cooling curves, the 

solidification time and evolution of the microstructure was 

recorded. The relation between the microstructure and 

the mechanical properties was also assessed by using quality 

index-strength charts developed for the alloy. This study 

shows that the microstructural features, especially the ironrich 

needles denoted as β-Al FeSi, and mechanical 

5 

properties are markedly affected by the different processing 

routes. The solidification rate exerts a significant effect on 

the coarseness of the microstructure and the intermetallic 

compounds that evolve during solidification, and this 

directly influences the tensile properties. 

KEYWORDS 

Aluminium cast alloys, gravity- and sand casting, 

gradient solidification technique, thermal analysis, 

porosity, quality index. 

Riassunto 

In questo lavoro sono state esaminate la microstruttura e le proprietà meccaniche di 

una delle leghe più comuni nell’industria delle leghe leggere, la lega Al-10%Si-0.4%Mg, 

colata per gravità sia in conchiglia che in sabbia. Le provette di trazione sono state 

ricavate da pala di ventilatore e sono state sezionate in tre posizioni caratterizzate da 

diverse velocità di raffreddamento. Inoltre è stato valutato il valore potenziale di resistenza 

a trazione ottenibile della lega, mediante produzione di campioni caratterizzati da diversa 

morfologia strutturale, basso contenuto di film ossidi microporosità, solidificati in 

ambiente controllato con metodologia di solidificazione a gradiente. Il comportamento 

in solidificazione è stato caratterizzato mediante analisi termica. Per mezzo di curve di 

raffreddamento, si è determinato il tempo di solidificazione e l’evoluzione della 

microstruttura. La relazione tra la microstruttura e le proprietà meccaniche è stata 

determinata usando specifiche carte di correlazione indice di qualità – resistenza. Questo 

studio dimostra che le caratteristiche microstrutturali, in particolare i grani aciculari ad 

alto tenore di ferro indicati come β-Al5FeSi, e le proprietà meccaniche sono 

marcatamente condizionate dai differenti percorsi preparativi. La velocità di 

raffreddamento esercita una significativa influenza sulla morfologia microstrutturale e 

sui componenti intermetallici che si sviluppano durante la solidificazione, influenzando 

così direttamente le proprietà di resistenza a trazione. 


INTRODUCTION 

The mechanical properties of cast aluminium alloys are very sensitive to 

composition, metallurgy and heat treatment, the casting process and the 

formation of defects during mould filling and solidification. The coarseness 

of the microstructure and the type of phases that evolve during solidification 

are fundamental in affecting the mechanical behaviour of the material. The 

mechanical properties obtained from gradient solidified samples are assumed 

to represent the upper performance limits of a particular alloy, due to the 

low level of defects obtained from the favourable melting and solidification 

procedure. 

Different processing routes offer different qualities and soundness of 

commercial components. The true potential of most alloys is seldom 

approached, and there exists a lack of data in the literature describing the 

variations in microstructure and related mechanical properties of alloy EN 

AC 43100. The present investigation was therefore carried out in order to 

Fig. 1: Fan blade cast in sand- and gravity moulds. 


elucidate and assess the influence of the casting 

process and resultant variations of microstructure 

on the mechanical properties, by extracting tensile 

specimens from Al-10%Si-0.4%Mg commercial cast 

components manufactured by two different casting 

methods; sand and gravity die casting. In addition, 

the inherent strength potential of this casting alloy 

will be investigated by manufacturing tensile test 

specimens using gradient solidification technology. 

The mechanical properties which have been 

measured in this work are: ultimate tensile strength, 

yield strength and fracture elongation. The 

microstructural features that have been evaluated 

are the secondary dendrite arm spacing (SDAS), 

pore fraction and precipitated phases such as the 

length of the iron-rich β-phase, Al 5 FeSi. 

MATERIALS AND 

EXPERIMENTS 

Cast component 

The components that have been studied are fan 

blades, as shown in figure 1, which are produced 

commercially in a great variety of shapes and sizes, 

and by different casting methods. This particular 

component is considered to be of interest to study 

and analyse due to the widespread use of different 

casting methods to manufacture fan blades, which 

in this case were manufactured by sand mouldand 

gravity die casting. 

The locations of the specimens extracted for study 

are shown in figure 1; from the top (T), middle 

(M) and bottom (B) parts of the fan blade. These 

locations were chosen due to their different 

solidification times (i.e. thicknesses), which results 

in different coarseness of the microstructure, 

intermetallic phases and defects. 

Cast alloy 

Table 1: Chemical composition of the casting alloy 

Composition Elements (wt. %) 

The commercial blades and specimens produced 

by gradient solidification experiments were cast 

using alloy EN AC 43100. The standard 

composition of the alloy and the actual 

composition are given in table 1. Strontium was 

Alloy Standard Si Fe Cu Mn Mg Zn Ti Cr Ni Al 

EN AC-43100 SS-EN 1706 9.0-11.0 0.55 0.10 0.45 0.20-0.45 0.10 0.15 - 0.05 Bal. 

EN AC-43100 Actual sample 9.96 0.47 0.11 0.25 0.35 0.097 0.018 0.01 0.01 Bal.

added in the form of Al-10%Sr master alloy at a 

level of 150- 200 ppm. 

Sand and gravity die casting 

The sand moulds used were made by hand. As a 

pattern, a fan blade cast in a gravity mould was 

used, which is as already mentioned, one of the 

methods normally used to manufacture this size 

and type of fan blade. The sand mould consisted 

of chemical bonded sand. 

Gradient solidification 

The same alloy as that used for the sand and gravity 

die cast components, shown in table 1, was cast 

into rods for further solidification studies 

in the gradient solidification equipment. The 

gradient solidification technique allows the 

production of a high quality material with a low 

content of oxide films, porosity and shrinkage 

related defects, and produces a homogenous and 

A) 

well-fed microstructure throughout the entire specimen. In this work a 

resistance-heated furnace with an electrically driven elevator was used, see 

figure 2a. Three different growth velocities, v, were used, 0.03 mm/s, 0.3 

mm/s and 3 mm/s, which correspond respectively to an SDAS of about 50, 

20 and 7 µm. At each velocity, three specimens were made. The specimens 

cast in the gradient solidification equipment were protected by argon in 

order to avoid oxidation. 

Thermal analysis 

In order to study the thermal history of the solidified components, 

thermocouples were inserted at three different positions in the sand cast 

fan blade, as indicated in the picture in figure 1. To record the cooling curves 

a data acquisition device from National Instruments was used. In each of 

the sand moulds three thermocouples were placed. The thermocouples 

were of S-type, which means that the different materials in the thermocouple 

wires are platinum and platinum-rhodium (90Pt-10Rh). When manufacturing 

the thermocouples, the wires are inserted in a two-hole Al 2 O 3 -tube to 

separate the wires, which are connected to a plug. The thermocouples 

were protected from the molten metal by a quartz glass tube. The maximum 

operating temperature for the S-type thermocouple is over 1500°C [1]. 

The recording rig is illustrated in figure 2b. 

B) 

Fig. 2: . Illustration of the instruments that have been used during the investigation. While a) demonstrates the gradient solidification equipment, 

b) presents the data acquisition rig that is connected to the sand moulds. 


RESULTS OF THE 

MICROSTRUCTURE AND 

MECHANICAL PROPERTIES 

INVESTIGATION 

Since specimens were extracted from fan blades 

with different thickness and at different sections 

it was necessary to machine them in order to 

obtain specimens suitable for tensile testing. The 

gradient solidified specimens have also been 

prepared in the same manner, and the dimensions 

of the samples used for tensile testing are shown 

in figure 3. In order to analyse the microstructure 

and measure the different microstructural features, 

a light microscope, image processing software and 

a scanning electron microscope, SEM, were used. 

Microstructure development 

The microstructure coarseness is defined as the secondary dendrite arm 

spacing, SDAS, which is also a function of the local solidification time, and 

quantifies the scale of the microstructure and its constituents. Examining 

the microstructure of specimens from both sand and gravity die casting, it 

has been observed that there is a wide scattering of SDAS and consequently 

variations of sizes and shapes of the intermetallic compounds within the 

same component. The needle-shaped β-phase, Al 5 FeSi also seems to coexist 

with the α-phase, Al 15 (Fe,Mn) 3 Si 2 , which has the morphology of Chinese 

script. As depicted in figure 4, the dendrites are equiaxed and the eutectic 

is finely modified. Comparing the microstructure formations, it can be noticed 

that in the case of gravity die casting the microstructure is finer, due to the 

higher cooling rate. Furthermore, the different coarsenesses of 

microstructure, the phases and their morphologies are comparable to the 

sand cast microstructure. 

It is relevant to note that even if Mn has been added to the alloy to act as 

A) B) 

Fig. 3: A schematic illustration of the tensile specimen (dimensions in mm). 

Fig. 4: . a) Illustration of the microstructure of a gravity die cast specimen while b) shows a sand cast microstructure. The black spot is a pore. 

 

 


 

 

 

an Fe-corrector and to promote the formation of 

Chinese script, see table 1, iron-rich Al 5 FeSi-needles 

are still formed. These needles act as local stress 

raisers in the matrix and are considered to be 

deleterious for the mechanical performance, which 

is why their formation and growth should be taken 

into consideration and as much as possible 

suppressed. 

In order to understand the local variations in 

microstructure, and to understand the formation 

sequences and development of constituent phases 

in the sand cast component, thermal analysis was 

implemented. When manufacturing the sand cast 

fan blades at the foundry, data for the cooling 

curves was recorded. In each of the sand moulds, 

three thermocouples were used, as shown in figure 1.

With this composition, see table 1, and casting 

technique, three different cooling curves have been 

obtained and are presented in figure 5a. The figure 

also shows that there exists clear variations in local 

solidification times which are directly related to 

the variations in component thickness, see also 

table 2. These variations in cooling time lead to 

major local differences in microstructures which 

affect the tensile performance of the cast 

components. The SDAS for the sand cast and 

gravity die cast component varied between 47-61 

µm and 20-45 µm respectively. 

Figure 5b shows that the freezing range is 

approximately between 595°C and 545°C. From 

the different cooling curves it is possible to see at 

which temperature and time the dendritic network 

starts to grow, at which temperature and time the 

eutectic starts to grow and also the total 

solidification time, from pouring the melt until the 

fraction solid has reached 100%. 

When the value of the derivative exceeds zero, 

figure 5b, the temperature is actually increasing 

due to the release of latent heat due to phase 

transformations. This is most evident in the case 

where a positive derivative is observed at about 

 

A) B) 

 

 

 

 

 

 

 

 

 

 

 

 

Mechanical properties 

The results of the tensile tests from the 

commercial fan blades and the laboratory 

prepared samples are shown in the diagrams in 

590°C when the dendritic network starts to grow. The next positive value 

of the derivative appears when the eutectic starts to grow, at about 570°C. 

When the temperature is around 550°C precipitations of a complex eutectic 

with intermetallics and a hardening phase such as Mg 2 Si starts to precipitate 

from the remaining melt. From the different thermocouples the following 

solidification times have been determined: 

Table 2. Solidification times for the different parts 

of the sand cast fan blade. 

Thermocouple Thinnest part Mid part Bottom 

(T) (M) part (B) 

Solidification time (s) 260 470 610 

Material thickness (mm) 12 20 28 

Examining the gradient solidified specimens, a significant discrepancy between 

the different solidification rates is observed. The higher the cooling rate 

the finer the microstructure, as illustrated in figure 6. The iron-rich 

β-phase Al 5 FeSi is found to be very short and very well dispersed for the 

two highest solidification velocities; 3 and 0.3 mm/s, with a corresponding 

SDAS ≈ 7 and 23 µm respectively, in comparison to the lowest velocity 

0.03 mm/s, with SDAS ≈ 47 µm, where long iron-bearing phases were found 

randomly distributed in the eutectic regions. 

Fig. 5: The thermal history of the sand cast component is registered in figure a) while in b) an analysis of the cooling curve 

obtained from the bottom section is performed. 

figure 7. The two different casting methods are compared in figure 7a, where 

each point represents a different position on the fan blade. A wide scatter 

in data is clearly observed but it can be noticed that the properties of the 

thinner parts seem to approach higher combinations of stress and strain 

values. 


In order to simulate the processes of the cast components, tensile test 

samples of the same alloy and corresponding microstructural coarseness 

have been produced by using gradient solidification. By producing a more 

homogenous microstructure with a lower level of defects, the material’s 

inherent potential was revealed, see figure 7b. The scatter in data in figure 

7b is ascribed to different levels of defects, and the higher stress-strain 

A) B) 

 

Fig. 6: Phases and microstructures for a) SDAS = 47 µm obtained at 0.03 mm/s and b) SDAS = 7 µm obtained at 3 mm/s. 

A) B) 

 

 

 

 

 

 

 

µ 

µ 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Fig. 7: The graph in a) illustrates the ultimate tensile strength and elongation to fracture for the gravity die and sand casting specimens where B is 

the bottom (thickest) section, M the middle and T the top (thinnest) part of the blade. The graph in b) presents the tensile properties for the 

gradient solidified specimens with a range of cooling rates. 

 

 

 

 

 

 

 


 

combinations are the limits for what this alloy 

might perform under ideal conditions. 

When quantifying the SDAS, five measurements 

were taken at three randomly chosen positions of 

the specimen. From these data a mean value was 

calculated. At high solidification rates the SDAS 

µ 

µ 

µ

appears to be very small and the microstructure 

has been found to be more uniform. As noticed, 

the overall effect of obtaining a small SDAS is an 

improvement in both ductility and tensile strength, 

see figure 8. 

When measuring the length of the iron-rich Al 5 FeSi 

needles (β-phase), the 15 longest needles observed 

in the microstructure were measured, considering 

the entire cross-section area of the sample, and a 

mean value was assessed. It should be mentioned 

that the needles were not chemically analyzed, but 

 

 

A) B) 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

µ 

µ 

µ 

µ 

µ 

 

µ 

 

 

were identified according to their morphologies and colours. A correlation 

was found between the length of the β-phase and the mechanical properties. 

An increase in needle length results in a reduction of tensile strength and 

elongation as can be seen in figure 9. Note that the iron-bearing β-phase 

was too short to be measured or not present at all in the specimen solidified 

at a rate of 3 mm/s, SDAS = 7 µm. It is also seen that even if these needles 

are relatively shorter in some cases than in corresponding samples, they 

exhibited lower properties. The premature fractures must therefore be 

due to other defects such as oxide film inclusions or other brittle and 

undesirable phases such as Al 8 Mg 3 FeSi 6 . These intermetallics are harmful to 

the mechanical properties since cracking of these phases reduces the ductility, 

 

Fig. 8: The graphs a) and b) illustrate the influence of SDAS on the tensile properties. 

A) B) 

 

 

 

 

 

 

 

µ 

µ 

µ 

µ 

 

 

 

 

µ 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 


 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

µ 

µ 

µ 

µ 

µ 

 

 

µ 

Fig. 9: The influence of the β-phase needle length on a) ultimate tensile strength and b) the elongation to fracture. 

 

 

 

 

 

 

 

 

µ 

 

 

 

µ 

µ 

µ 

µ

µ 

 

 

 

 

 

 

 

 

µ 

µ 

µ 

µ 

 

 

 

 

 

 

 

 

µ 

Fig. 10: This graph shows the relation between the β-phase length and the SDAS. 

and their formation reduces the Mg available in the melt to form Mg 2 Si 

which has a hardening effect and improves both the yield and ultimate 

tensile strengths. 

Generally, larger SDAS is associated with lower cooling rates and also with 

longer β-phase needles. The local solidification time influences the growth 

of the intermetallics that develop and grow during the solidification process, 

and hence a relation between SDAS and the β-phase length is observed as 

shown in figure 10. 

When a tensile sample contains pores it is reasonable to assume that the 

region of porosity will yield first due to the reduced load bearing area 

concentrating the stress near the voids. The area fraction of porosity was 

determined by measuring the area just below and at the fracture surface. 

 

A) B) 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

µ 

µ 

µ 

µ 

µ 

 

 

 

 

 

 

 

Fig. 11: The graphs a) and b) illustrate the influence of porosity on the tensile properties. 

 


 

 

 

 

 

 

 

 

The first mentioned was measured by an image 

analyzer and the latter was photographed and 

observed by scanning electron microscopy, SEM. 

This method enables an examination of the 

distribution and size of the pores. 

In the present study, it has been observed that the 

cooling rate, SDAS in this case, seem to govern 

the length as well as the percentage area fraction 

of porosity to some extent, see figure 11. It appears 

that a reduction in SDAS results in smaller average 

pore size and a reduced area fraction of porosity. 

It is obvious that as the solidification time is low, 

less time will therefore be available for the diffusion 

of the hydrogen into the interdendritic regions 

which results in small sizes and fraction of pores. 

But comparing the porosity formation, the gradient 

solidified samples with SDAS ~ 23 µm are 

associated with largest percentage area fraction 

of porosity which might be due to the arrestment 

of premature pores as they become entrapped by 

the advancing solidification front. 

The discrepancy could be derived from the theory 

that involves the harmfulness of the oxide films, 

which suggests that when an oxide particle is 

approached by the solidification front, the particle 

experiences the hydrogen-rich environment 

produced by the rejection of gas from the 

advancing solid. Furthermore, the access of gas by 

diffusion into the air pocket in the gap of the oxide 

particles, the pore will start to form and grow. As 

the freezing rate is slow and the particle is poorly 

wetted by the melt, time will therefore be available 

for more hydrogen to diffuse resulting in pore 

expansion and growth. 

 

 

 

 

 

 

 

µ 

µ 

µ 

µ 

µ

In the case of the larger SDAS ~ 47 µm, it is 

therefore reasonable to assume that due the slow 

cooling conditions many air pockets formed within 

the liquid or due to interaction with oxide particles 

have been driven in front of the solidified front 

without being engulfed and in that case not been 

detected in the gradient solidified specimens due 

to the solidification mode. Another reason might 

be the longer time available for the hydrogen to 

diffuse and move in front of solid-liquid interface 

out of the sample into the surrounding 

environment. 

The scatter in the data is too great to derive any 

clear correlation between porosity level and the 

tensile strength and ultimate elongation, see figure 

11 a and b. According to literature, it is stated [2] 

that the static tensile properties such as ultimate 

tensile strength, yield strength and elongation to 

A) B) 

DISCUSSION 

Generally, the mechanical properties of metals are 

widely influenced by their microstructures. When 

defining the microstructure many parameters have 

to be taken into account, and these include among 

others the secondary dendrite arm spacing, the 

size and shape of precipitated phases such as silicon 

and iron-bearing phases, grain size, porosity, etc. 

In this investigation, many of the parameters 

mentioned above have been measured and 

correlated to the mechanical properties. 

Constituents that may have deleterious effect on 

the mechanical properties are defects such as 

oxide films or undesired phases and inclusions. 

A comparison of the results of the tensile testing 

with the casting method and the corresponding 

fracture, were all decreased with an increased degree of porosity. On the 

contrary, according to [3-7], the reduction in tensile properties has almost 

no correlation with the average volume fraction of porosity. In fact, the 

decrease in tensile properties was attributed to the length and/or the area 

fraction of defects in the fracture surface. Development of a high fraction 

of porosity may also be due to Fe in the melt. The long, needle-shaped iron 

intermetallic phase formed is expected to cause severe feeding difficulties 

during solidification. The morphology of the β-phase blocks the interdendritic 

flow channels, which is why it is proposed that higher iron contents in the 

alloy are associated with higher levels of porosity. 

SEM examination was also performed in order to study what kinds of pores 

are frequently found on the fracture surfaces in this work. Figure 12a 

illustrates a pore seen in a sand cast specimen and 12b is seen in a gradient 

solidified specimen at a solidification rate of 3 mm/s. 

Worth to indicate is that the melt hydrogen content has not been measured 

and assumed in this case to be constant for all the alloys since they have 

been produced under similar conditions. 

Fig. 12: Illustration of pores. a) Showing sphericity of a possible gas pore b) The cavity is observed at the edge of the 

fractured surface shows a possible gas-shrinkage pore. 

specimens with three different solidification rates obtained by gradient 

solidification shows clearly that the tensile behaviour and production method 

is related by three different stress-strain curves. The material exhibiting 

SDAS ≈ 7 µm results in the highest stress-strain combinations as shown in 

the diagram in figure 13. The main reason for this behaviour is the finer 

microstructure, a more homogeneous distribution of the intermetallic phases 

and lower levels of porosity. Even though there are no specimens with 

corresponding SDAS extracted from commercial components available in 

this study, this data clearly shows the inherent potential strength of the 

alloy. 

The stress-strain values for the sand cast specimens follow the curve for 

SDAS = 47 µm (obtained by gradient solidification), but fracture occurs at 

lower stress and strain levels, see figure 13. The stress-strain curve for 

SDAS = 23 µm (by gradient solidification) corresponds well to the stressstrain 

values of the gravity die cast specimens. 

The reason why the sand and gravity die cast specimens fractured much 

earlier than the gradient solidified specimens, and exhibited different 


µ 

µ 

µ 

µ 

µ 

µ 

µ 

µ 

 

 

 

 

 

 

Fig. 13: This graph shows tensile test curves for the gradient solidified specimens and 

includes mechanical properties data for the gravity die and sand cast specimens. 

combinations of strength and fracture elongation values may be due to the 

formation of brittle and coarse intermetallics such as the iron-bearing 

β-phase and an inhomogeneous microstructure. Even if similar SDAS has 

been obtained from the gradient solidification equipment, the variations in 

the β-phase needles, local coarsenesses of microstructures and sizes of 

intermetallics are likely to be different due to microsegregation in the 

castings. Segregation of elements is likely to occur in castings and this will 

affect the development and solidification sequences of the microstructure 

and lead to local variations in mechanical properties. At the thinner part of 

the fan blades the β-phase needles are likely to be short and have no time 

to grow and coarsen as is the case in thicker parts where the local iron 

 

A) B) 

 

 

 

 

 

 

 

 

µ 

µ 

 

 

 

 

 

 

 

 

 

 

 

 

 


level might also be different. It has also been shown 

that a finer microstructure with smaller SDAS is 

beneficial for the mechanical properties. A finer 

microstructure means larger areas of grain 

boundary leading to more difficulties for 

continuation of slip and dislocation movement 

during deformation, hence yielding stronger alloys. 

A simple tool that might be used to predict the 

impact of changes in the chemical composition, 

solidification rates and routes, microstructure and 

heat treatment on the tensile performance is by 

implementing a quality index, Q. Its practical value 

originates from the fact that the mechanical 

“quality” of an alloy is expressed by a single 

numerical parameter, whose value can be read off 

a quality index chart or calculated through a simple 

equation relating the elongation to fracture, s , the f 

ultimate tensile strength, UTS, and d, which is an 

empirical parameter chosen to be around 150 MPa 

for EN AC 43100 (Al-7%Si-0.4%Mg) to make Q 

more or less independent of the yield strength of 

the alloy, as seen in equations 1-2, [8-10]: 

Q = UTS + d * log (100*s ) (1) 

f 

n 

σ = Kε 

(2) 

where σ is the true flow stress, K is the alloy’s 

strength coefficient, ε is the true plastic strain and 

n the strain hardening exponent. 

In order to produce the charts in figure 14, a mean 

value for the strength coefficient, K, has been 

calculated. For the commercial cast components, 

figure 14a, K is approximately 410 MPa and for 

the gradient solidified materials K is around 430 

MPa, figure 14b. The strain hardening exponent, n, 

is numerically equal to the necking onset strain 

Fig. 14: A quality index chart for the alloy studied obtained for gravity- and sand cast samples is presented in a) while b) represents the gradient 

solidified specimens. The mean value of the strength coefficient is a little higher for the gradient solidified specimens. 

 

 

 

 

 

 

 

 

µ 

µ 

µ

and can be used to define the relative ductility 

parameter, q, see equation 3. 

q = s f /n (3) 

Defining q = 1 means the onset of necking 

represents samples that fail either at the onset of 

necking or beyond, while q < 1 identifies lower 

quality or less ductile samples. For instance q = 

0.5, 0.2, 0.1 identifies samples that fail at 50%, 20% 

and 10% of their maximum uniform strain, 

respectively. 

As observed, the points corresponding to 

SDAS = 52 µm are located near the bottom of 

the chart, which may be due to their coarser 

CONCLUSIONS 

As an outcome from the current investigation, the 

dendritic cell size, SDAS, and the length of ironbearing 

phase whether as needles or Chinese 

script, governs the tensile strength and ductility of 

Al-Si-Mg alloys. The early fracture of sand and 

gravity die cast specimen can be related to the 

length of the iron-bearing phase, oxide films 

entrapment, and inhomogeneous microstructure 

and to the Si particle size and morphology, which 

determine the rate of particle cracking with the 

applied strain. 

• A finer microstructure, small and well 

distributed iron-rich compounds and low levels 

of porosity exert a beneficial impact on the 

resulting mechanical properties. 

• The locally increased levels of iron due to 

ACKNOWLEDGEMENTS 

REFERENCES 

1. Metals Handbook; ASM International, 1998, pp. 

663-665. 

2. M. Avalle, G. Belingardi, M.P. Cavatorta and R. 

Doglione, International Journal of Fatigue Vol. 24 

, (2002), pp. 1-9. 

3. C.H. Cáceres and B.I. Selling , Materials Science 

and Engineering A, Vol. 220, Issues 1-2, (1996), 

pp. 109-116. 

4. M. Harada, T. Suzuki and I. Fukui , Journal of the 

Japan foundrymen´s society Vol. 55, No. 12, 

(1983), pp. 47-50. 

microstructure and the length of the β-phase. As the strength of the gravity 

die cast specimens increases, the data points shift toward the upper right 

corner. 

Regarding the gradient solidified samples, the data points for the 

SDAS = 7 µm specimens, with the finer microstructure, exhibit the highest 

values of Q and q. These values seem to be the only ones that reach the 

limit imposed by the necking line q = 1. 

As the SDAS and the level of deleterious microstructural compounds are 

decreased, the strength and ductility of the alloy is increased and so does 

the Q, allowing for a possible comparison between different alloys and 

processing routes. Furthermore, the Si particles, playing a role as reinforcing 

phases, decrease in size and the shape becomes more fibrous, which has an 

overall beneficial effect on the mechanical properties. 

possible segregations might have resulted in longer and larger fractions 

of β-phase needles. The main reason for the remarkable reduction of 

the mechanical performance of commercial components is proposed 

to be due to the presence of the locally induced stresses by the needles 

which break, link and propagate the cracks in an unstable manner. The 

higher the solidification rate the shorter the length of the needles and 

the sounder the material. 

• The gradient solidification method shows that the EN AC 43100 alloy 

used for commercial cast components processed by sand and gravity 

die casting, has an improvement potential to provide sounder castings 

with substantial ductility. 

• The addition of Mn did not alter the morphology of iron-bearing βphase 

into a more compact Chinese script. Instead, these iron 

intermetallics seem to be coexisting in the microstructure. 

• Porosity is not found to be the primary parameter controlling the tensile 

strength and ductility, but it is worth bearing in mind that the reduction 

of the load bearing area generally has a harmful influence on the overall 

properties. 

The authors acknowledge Johan and Jan Tegnemo at Mönsterås Metall AB for their invaluable help and support for making the 

experiments in the foundry and Nordic Industrial Fund for the financial support within the Nordic Project STALPK. 

5. M.K. Surappa, E. Blank and J.C. Jaquet, Scripta Metallurgica, Vol. 20, No. 9, 

(1986), pp. 1281-1286. 

6. J.G. Conley, J. Huang, J. Asada and K. Akiba, One Hundred Third Annual 

Meeting of the American Foundrymen’s Society; Rosemont, IL; USA, (1998). 

pp. 737-742. 

7. C.H . Cáceres, Scripta Materialia Vol. 32 No.11, (1995), pp. 1851-1856. 

8. C.H. Cáceres, I.L. Svensson, J.A. Taylor, International Journal of Cast Metals 

Research, Vol. 15, No. 5, (2003), pp. 531-543. 

9. C.H. Cáceres, International Journal of Cast Metals Research, Vol.12, (2000), 

pp. 367-375. 

10. M. Drouzy, S. Jacob and M.Richard, AFS International Cast Metals Journal, 

Vol. 5, No. 2, (1980), pp. 43-50. 


COMPARATIVE STUDY OF HIGH 

TEMPERATURE WORKABILITY OF ZM21 

AND AZ31 MAGNESIUM ALLOYS 

1 M. El Mehtedi, 1 L. Balloni, 1 S. Spigarelli, 1 E. Evangelista, 2 G. Rosen, 3 B.H.Lee, C.S.Lee 

1 Dipartimento di Meccanica, Università Politecnica delle Marche, Ancona, Italy 

2 ALUBIN, Kiryat Bialik, Israel 

3 Department of Materials Science and Engineering, Pohang University, Pohang, Korea 

Abstract 

High temperature regime, 300-450°C for Mg-Al-Zn alloys, is currently used in primary 

processing, such as rolling and extrusion, as well as for secondary operation like forging. 

The knowledge of temperature and strain rate proper combination (processing window) 

as well as the microstructure evolution occurring during hot deformation clarifies the 

relationships between forming variables and final properties of components. Numerous 

data on AZ31 and few other Mg-Al alloys, produced by laboratory testing, are available 

in the scientific and technical literature. The ZM21, Mg-2Zn-1Mn, by contrast, is 

characterized by absolute lack of scientific data. In the alloy the addition of manganese, 

by suppressing the formation of beta phase, increases the solidus temperature that 

results in the larger processing window than in AZ31. The benefit requires extensive 

analysis aimed at optimizing the deformation variables that affect the microstructure 

refinement under dynamic and static recrystallization. The high-temperature plastic 

deformation and the microstructure evolution of the ZM21 were thus investigated in 

the temperature range between 200 and 500°C and results were analysed and compared 

with those of a conventional heat-treated AZ31. 


Riassunto 

Il regime delle alte temperature 300-450°C, è comunemente 

usato per le operazioni primarie di deformazione plastica, 

come la laminazione e l’estrusione, come pure per quelle 

secondarie come la forgiatura o lo stampaggio delle leghe 

Mg-Al-Zn. La conoscenza della combinazione adatta 

temperatura-velocità di deformazione (finestra di processo) 

e della evoluzione della microstruttura fa luce sulle relazioni 

tra variabili di deformazione e proprietà dei componenti. 

Numerosi dati, ottenuti con prove di compressione o 

torsione, sono disponibili sulla lega AZ31 e su pochi altre 

leghe della classe Mg-Al. Sulla lega ZM21 (Mg-2Zn-1Mn) 

non sono invece disponibili dati scientifici. Nella lega 

l’addizione del Mn sopprime la formazione della fase beta 

ed aumenta la temperatura del solidus, producendo una 

finestra di processo più ampia di quella dell’AZ31. Questi 

aspetti richiedono analisi approfondite sono rivolte 

all’ottimizzazione delle variabili di deformazione che poi 

governano l’affinamento della microstruttura per effetto 

della ricristallizzazione sia dinamica che statica. La 

deformazione plastica ad alta temperatura e l’evoluzione 

della microstruttura della lega ZM21 sono state analizzate 

nell’intervallo di temperatura 200-500°C e i risultati sono 

confrontati con quelli di una lega convenzionale trattata 

termicamente.

1. INTRODUCTION 

Lightweight constructions are becoming more and 

more important for the automotive industries due 

to increase of fuel prices and legislative 

requirements like CAFÉ in the US, and the directive 

or the control of CO 2 emissions in the EU [1,2]. 

Magnesium alloys for their unique combination of 

light weight, high specific strength and stiffness and 

high recycling capability are candidate to replace 

steel, iron and, in some cases, also aluminum parts. 

The growing interest in magnesium alloys is 

expected to increase in the next years [3]. The 

use of components manufactured via die-, sand-, 

mould- and rheo-casting is fairly well introduced 

in a wide number of applications which are ranging 

from steering wheels, instrumental panels, gear box 

housing and even to hybrid engine blocks. The 

scenario for the automotive industry is that the 

use of magnesium alloys for many components, 

mainly as castings, is further expanding in the short 

term, whereas the application in body and chassis 

components, including sheet and extrusions, is 

anticipated in the medium term [4]. 

A driving force for the introduction of wrought 

alloys is their property enhancement, since casting 

alloys are affected by porosity, which would be of 

particular importance for structural and crashrelevant 

parts. Initial applications for this category 

are expected for extrusions and forgings, to be 

followed by sheets. These requirements may 

address researches towards the modification of 

known alloys as well as the development of new 

ones aiming at improving both mechanical 

performances, also by ageing (T6), and high 

temperature stability. The growing interest in use 

of magnesium alloys is expected to increase in the 

next years, and to be sustained by research [5]. 

From recent magnesium alloys conferences [6], the 

following items stand out as limiting or retarding 

factors for the use of wrought alloys: 1) poor cold 

forming properties; 2) limited number of wrought 

alloys and lack of processing data. As far as the 

item 1) is concerned, it is important to note that 

magnesium alloys are difficult to work at low 

temperature (

EXPERIMENTAL 

The AZ31 experimental result data belong to the authors’ wide database 

obtained in the last 3 years by testing a batch of AZ31 with different initial 

microstructure: ingot, extruded, heat treated, overaged. These investigations 

aimed at clarifying the microstructure features that significantly affect the 

high temperature deformation of the alloy; the data were partly published 

[11,12] and used to identify the processing window (strain rate and 

temperature conditions) that is successfully used for industrial forming 

operations. Warm forming experiments were also carried out to assess the 

AZ31 sheet formability [13]. The present data were obtained by testing 

the specimens machined from rolled AZ31 heat treated at 385°C for 14 h, 

at the Pohang University. 

The ZM21 alloy was produced and extruded by Alubin, Israel. The specimens 

for torsion testing, with gauge radius R, and length L, 5 and 10 mm respectively, 

were machined from extruded rods. Torsion tests were carried out in air 

on a computer-controlled torsion machine, under equivalent strain rates 

ranging from 10 -3 to 5 s -1 and temperatures from 200°C to 400°C. Specimens 

were heated 1°C/s by induction coils and maintained 5 minutes before 

torsion to stabilize the testing temperature that was measured by 

thermocouple in contact with the gauge section. Specimens just after the 

fracture were rapidly quenched with water jets to avoid microstructure 

RESULTS AND DISCUSSION 

The initial microstructure of the alloys before testing is displayed in Figure 

1. The heat-treated AZ31 exhibited an equiaxed and fine microstructure. 

The microstructure of the extruded ZM21 is composed by duplex grain 

size, the small and the large grains were 20 and 60 µm respectively. 

Figures 2 and 3 show typical equivalent stress vs. equivalent strain curves 

obtained by testing at 200 and 400°C the rolled and heat treated AZ31 and 

the extruded ZM21, respectively. The flow curves shape of the rolled AZ31 

does not differ appreciably from that observed in the cast [11,12]. At 

A) B) 


modifications during slow cooling. 

The von Mises equivalent stress, σ, and equivalent 

strain, ε, were calculated using the relationships: 

σ = 

3 M 

( 3 + m'+ 

n' 

) 

2 π R 

3 

2 π N R 

ε = 

2) 

3 L 

where N is the number of revolutions, M is the 

torque, m’ (strain rate sensitivity coefficient) at 

constant strain is ∂ log M / ∂ log N, 

• 

and n’(strain 

hardening coefficient) at constant strain rate is 

∂ log M / ∂ log N; at the peak stress, clearly n’=0. 

To reveal the microstructure, longitudinal sections 

at the periphery of the sample gauge, i.e. in the 

point where the strain and strain rates assumed 

the values calculated by equations 1) and 2), were 

investigated by means of optical microscopy. 

200°C, under all strain rates, the flow curves of 

both alloys exhibit well defined σ y , ~60 MPa for 

AZ31 and ~20 MPa for ZM21, a continuous 

increase of work hardening up to σ max , which is 

larger in AZ31, followed by fracture that in ZM21 

occurs at larger strain. At 400°C, AZ31 the flow 

curve, after yielding and very limited workhardening, 

gets a steady state region, typical 

behaviour of dynamic recovery (DRV), up to 

fracture. The ZM21 flow curves display yielding 

Fig. 1: Microstructure of the rolled and heat-treated AZ31 (a) and of the extruded ZM21 (b). 

1)

and rapid hardening up to σ peaks located at ε~0.7, 

irrespective of the different strain rates, followed 

by a continuous softening towards the fracture 

that occurs at e larger than in AZ31; at 400°C and 

strain rate 0.05 s -1 the maximum strain value is 

2.5. The flow curve exhibiting a peak followed by 

softening in stationary state is indicative of dynamic 

recrystallization (DRX) that nucleates in proximity 

of the peak producing new stable grains whose 

size controls the softening and the σ of stationary 

state. 

In the present instance, the absence of stationary 

state and the decrease of σ up to fracture may be 

indicative of three processes promoting the 

continuous softening: (i) the onset of DRX around 

the peak, producing new small grains, (ii) the fast 

new-grain growth at the critical size for (iii) crack 

nucleation and growth. 

A significant decrease in the strain to failure in 

torsion can be observed (Figure 4), in particular 

in the high-temperature region; an additional 

difference is the observation of a well defined 

yielding in ZM21, (Figure 3) i.e. of the presence of 

a short strain range without load increase, an effect 

A) 

 

B) 

 

 

 

 

 

 

 

 

 

 

 

 

A) 

B) 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Fig. 3: Equivalent stress vs. equivalent strain curves obtained in torsion for 

ZM21 at 200 (a) and 400°C (b). 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 


Fig. 2: Equivalent stress vs. equivalent strain curves obtained in 

torsion for AZ31 at 200 (a) and 400°C (b). 

probably due to a sudden increase in the number of 

dislocations unpinned from their solute atmosphere 

(either Zn or Mn). 

The temperature and strain rate dependence of the 

peak stress are shown in Figure 5. The experimental 

data are well described by the equation: 

ε& = A[ sinh( 

ασ) 

] n 

exp( −Q 

/ RT) 

3) 

where A and α are material parameters, n is a 

constant, R is the gas constant, T is the absolute 

temperature, and Q is the apparent activation energy 

for high-temperature deformation process. In the 

AZ31 a best-fitting procedure gave α=0.017 MPa-1 and n= 4.4 for temperatures between 200°C and 

400°C. The activation energy, calculated by plotting 

sinh(as) as a function of 1/T, was Q=140 kJ/mol. The 

value is relatively close to the one of self-diffusion of 

Mg (135 kJ/mol) or for diffusion of Al in Mg (143 kJ/ 

mol). In the case of the ZM21, α was close to 0.049 

MPa-1 , and n ranged between 2.6 and 4.4 (average 

n=3.3). The activation energy for high temperature 

deformation was close to 200 kJ/mol, higher than in 

the case of AZ31.

A) B) 

 

 

 

 

 

 

 

 

 

 

 

 

 

Fig. 4: Equivalent strain to failure for AZ31 (a) and ZM21 (b). 

A) B) 

 

 

 

 

 

 

 

ασ 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Fig. 6: Zener-Hollomon parameter as a function of peak-flow stress for AZ31 and 

ZM21. 


 

 

 

 

 

 

 

 

Fig. 5: Peak flow stress dependence on applied stress as a function of temperature for AZ31 (a) and ZM21 (b). 

 

 

 

 

ασ 

α 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

ασ 

Figure 6 plots the Zener-Hollomon parameter, Z, 

defined as 

[ ( ) ] n 

sinh 

 

 

 

 

 

Z = ε& 

exp( Q / RT) 

= ασ 

4) 

as a function of peak stress, with Q= 170 kJ/mol 

and α=0.034 MPa-1 (mean values of those obtained 

experimentally for AZ31 and ZM21). Analysis of 

Figure 6 clearly demonstrates that the peak flow 

stresses are higher in the AZ31 than in the ZM21. 

Figure 7 shows typical microstructure of the AZ31 

quenched after torsion testing. At 200°C, and a 

strain rate of 10-2 s-1 , the structure is mostly 

composed by slightly elongated grains (Figure 7a), 

with very fine recrystallized grains on several grain 

boundaries. At 300°C, under the same strain rate, 

the recrystallized fraction is substantially higher 

than at the lower temperature, and the average 

size of the recrystallized grains is larger (Figure

A) B) 

C) D) 

Fig. 7: Microstructure of the AZ31 tested at: 200°C-10-2 s-1 (a) 300°C-10-2 s-1 (b) and 400°C- 1 s-1 (c); (d) 400°C-10-3 s-1 . 

7b). Fully recrystallization and localised grain 

growth at 400°C, in the high strain rate range, 

resulted in the presence of some large grains 

heavily deformed by twinning, albeit a relatively 

large fraction of the structure is composed by fine 

equiaxed grains. Under lower strain rate, the 

structure is composed by a homogeneous 

distribution of relatively coarse equiaxed grains 

(Figure 7d). 

The mechanism of dynamic recrystallization (DRX) 

has been analysed by McQueen and Konopleva 

[15]; twinning in fact takes place before the other 

mechanisms, except basal glide that occurs in 

favourably oriented grains. At low strains, twinning 

favourably reorients grains for slip. As deformation 

reaches a sufficiently high value, DRX starts where 

high misorientations have been created by 

accumulation of dislocations, i.e. where slip has 

occurred on several slip system (near grain 

boundaries and twins). The new fine grains form a 

mantle (or a “necklace”) along grain boundaries 

and deform more easily than the grain core, thus repeatedly undergoing 

recrystallization [12]. 

Figure 8 shows typical micrographs of the microstructure of ZM21 tested 

in torsion. The structure, even at 300°C, is largely unrecrystallized, with 

elongated grains; at 400°C, the grains are mostly equiaxed, but elongated 

structures still persist. The strong decrease in flow stress after the peak 

should be attributed to the late onset (at strain close to ε=0.7-0.8) of 

recrystallization and coarsening phenomena, that in this alloy exhibit quite 

different kinetics than in the case of the AZ31. Only at the highest 

temperature (500°C), the microstructure underwent complete 

recrystallization and grain growth, a process that resulted in the presence 

of coarse (50 µm) grains. 

The above analysis gives some interesting preliminary indications on the 

high-temperature formability of the extruded ZM21 alloy. In particular: 

1. the intrinsic ductility, i.e. the equivalent failure strain in torsion, is lower 

in ZM21 than in AZ31; 

2. the fraction of recrystallized structure, at all the investigated 

temperatures, is lower in ZM21 than in AZ31; 

3. the maximum equivalent stress, i.e. the working load, is higher in AZ31 

than in ZM21; yielding has been observed in ZM21, thus suggesting a 

solute strengthening role played by either Zn or Mn; 


A) 

C) 

Fig. 8: Microstructure of the ZM21 tested at: 300°C-5 s-1 (a) 400°C- 5x10-2 s-1 (b) and 500°C- 5 s-1 (c). 

4. the different composition of the ZM21 enables the use of higher 

working temperature, up to 500°C, i.e. well above the upper limit 

allowed for AZ31. On the other hand, even though strain to failure is 

relatively high at 500°C, the microstructure undergoes extensive grain 

CONCLUSIONS 

The high temperature workability of the extruded ZM21 alloy has been 

investigated by torsion tests between 200 and 500°C. Both mechanical and 

microstructural results were compared with those observed in a rolled 

AZ31 alloy; the comparison indicated that the strain to fracture in torsion 

were significantly lower in ZM21 than in AZ31, even though in general in 

the former alloy the ductility exceeded the minimum allowable value of 

ε=1. These differences in mechanical response were attributed to a lower 

tendency to dynamic recrystallization; while in rolled AZ31 early dynamic 

recrystallization was observed at temperatures as low as 200°C, 

microstructural observations clearly suggested that in the extruded ZM21 

the lower limit for DRX was shifted toward significantly higher temperatures. 

B) 


growth, that precludes the fine grain size that 

has been observed to be a prerequisite for 

optimum sheet formability. 

ACKNOWLEDGEMENTS 

The authors greatly acknowledge the Ministries 

of Foreign Affairs of Italy and Korea who in the 

frame of “2004-06 Joint Scientific and Technological 

Research Program” promoted the research by 

funding the researchers’ mobility. 

The valid support of D. Ciccarelli in experimental 

activity has been appreciated.

REFERENCES 

1) CAFÉ, http:/www.nhtsa.dot.gov 

2) Communication from the Commission to the 

Council and European Parliament: 

Implementing the Community strategy to 

reduce CO emissions from cars. Third annual 

2 

report (reporting year2001), COM (2002) 

693. 

3) EU Directive on the End of Life of Vehicles, 

2000/53/EC. 

4) S. Schumann, Mat. Sci. Forum 488-489 (2005), 

1-8. 

5) Ya Zhang, X. Zeng, C. Lu, W. Ding, Y. Zhu, Mat. 

Sci. Forum, 488-489,(2005), 123. 

6) Magnesium Technology 2006, A.A. Luo et al, 

eds. TMS, 2006. 

7) F.J. Humphreys, M. Hatherly: Recrystallization 

and related annealing phenomena. Pergamon 

Press, Oxford, 1996. 

8) J. Koike, Mat. Sci.Forum, 419-422 (2003), 199. 

9) H. Utsunomiya, T. Sakai, S. Minamiguchi, H. Koh, Magnesium Technology 

2006, A.A. Luo et al. eds. TMS, 2006, 201. 

10) J. Bohlen, J. Swoistek, W.H. Sillekens, P.J. Vet, D. Letzig, K.U. Kainer, 

Magnesium Technology 2005, TMS, 2005, 241. 

11) S. Spigarelli, M. El Mehtedi, E. Evangelista, J. Kaneko, Metall.Sci.Techn. 23 

(2005) 11-17. 

12) S. Spigarelli, M. El Mehtedi, M. Cabibbo, E. Evangelista, J. Kaneko, A. Jäger, 

V. Gartnerova, mater. Sci.Eng. A 462 (2007), 197. 

13) A. Forcellese, M. El Mehtedi, M. Simoncini, S. Spigarelli, Key Engineering 

Materials 344 (2007) 31. 

14) S. Spigarelli, E. Evangelista, M. El Mehtedi, L. Balloni, Mechanical And 

Microstructural Aspects Of High Temperature Formability Of AZ31 

Sheets, Proceeding 10 ESAFORM 2007, Zaragoza, Spain, 1305. 

15) H.J. McQueen, E.V. Konopleva, in Magnesium Technology 2001, J. Hryn 

ed., TMS, 2001, p. 227. 


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