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Wilm noted that an aluminium alloy he was working on had the remarkable
ability to harden slowly at room temperature after having previously been
quenched from a temperature just below its melting point. The effect was
observed, quite fortuitously, while Wilm, working on a contract from the
Prussian government, was attempting to develop alloys for cartridge cases
which were lighter than the 70/30 brass which was normally employed.
Between 1909 and 1911 Wilm filed several Patent Applications, and
assigned his rights in this ‘age-hardening’ invention to the Durener-
Metallwerke in Duren, who subsequently marketed the alloys under the trade
name Duralumin. During the First World War large quantities of age-hardened
aluminium alloys were used by the combatants, first for Zeppelins and then for
other types of aircraft. It was found that the strengthening reaction could be
slowed down, very significantly, by refrigeration. This made it possible to
quench aluminium alloy rivets, and store them in the soft condition in a
refrigerator. The age hardening process began only after the rivet head had
been closed after insertion into the aircraft structure.
A great deal of the metallurgical research stimulated by the introduction of
Duralumin was co-ordinated in Great Britain from 1911 by the Alloys
Research Committee of the Institution of Mechanical Engineers. The 11th
report of this committee, published in 1921, summarized eight years of work
by the National Physical Laboratory (NPL). The outstanding result of this
work was the development of ‘Y’ alloy (4 per cent Cu, 2 per cent Ni, and 1.5
per cent Mg), which retained its strength to moderately high temperatures and
was extensively used for pistons and cylinder heads. Y alloy was originally
used in the cast form. High Duty Alloys was established in 1928 by Colonel
W.C. Devereux to produce variants of Y alloy in the forged condition for
aircraft engine use. From these experiments stemmed the RR series of light
alloys, jointly developed by Rolls Royce and High Duty Alloys.
The first rational explanation of the age hardening process in light alloys

such as Duralumin and Y alloy was provided in 1919 by Paul Merica and his
colleagues Waltenberg and Scott at the National Bureau of Standards in
Washington. They found that when an alloy such as Duralumin was heated to
temperatures close to its melting point most of the alloying constituents were
taken into solution by the matrix. Quenching retained the dissolved metals in
this supersaturated solid solution which was, however, somewhat unstable at
room temperature: minute crystals of various intermetallic compounds were
slowly precipitated. Provided these crystals were below a certain critical size,
invisible to the optical microscope, they strained and distorted the aluminium
lattice and acted as mechanical keys, which inhibited plastic flow in the alloy
and increased its strength. The aluminium alloys which behaved in this way
were unique only in the sense that the precipitation effects manifested
themselves, quite fortuitously, at room temperature. A wide range of other
alloys were soon identified in which precipitation could be induced by ageing
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at elevated temperatures and new precipitation hardening alloys of all types
were rapidly developed. In 1934 Dr Maurice Cook of ICI was able to review
the precipitation hardening characteristics of 37 copper alloy systems.
Beryllium and beryllium alloys
Beryllium, first identified as an element by Wöhler in 1828, was at first
regarded merely as a chemical curiosity. Soon after the introduction of the
incandescent gas mantle by Auer von Welsbach in 1885 it was found that the
strength and durability of the ash skeleton could be greatly improved by
adding small quantities of beryllium nitrate to the mixtures of thorium and
cerium nitrates used to impregnate the fabric of the mantle. This treatment was
probably the major outlet for beryllium until the end of the First World War.
Attempts to produce beryllium by the electrolysis of a fused bath of
beryllium chloride were first made in 1895 by Wilhelm Borchers. Goldschmidt
in 1915 found that fused fluoride baths offered better production prospects.

This approach, refined by Stock, Praetorius and Priess, yielded relatively pure
beryllium for the first time in 1921. The properties of the metal obtained
confirmed theoretical predictions that beryllium would be a light metal with a
density around 1.8 grams per cm
3
and that 1ts direct modulus of elasticity
would be significantly higher than that of steel. If its ductility could be
improved, it was likely to be the light, stiff metal which had long been sought
by the aircraft industry.
The alloying behaviour of beryllium was studied by Regelsberger in 1924.
Additions of up to 10 per cent by weight of beryllium improved the hardness
and tensile strength of magnesium without making it brittle. Up to 10 per cent
of beryllium could also be dissolved in copper to produce a pale yellow alloy.
The colour of the alloy improved as the beryllium content decreased, and
Regelsberger was the first to mention the beautiful golden colour of copper
containing around 2 per cent of beryllium.
Research on beryllium in the United Kingdom was initiated by the Imperial
Mineral Resources Bureau at the Metallurgy Department of the NPL in 1923.
By 1926 Dr A.C.Vivian had produced solid deposits of metal by the
electrolysis of fused baths of beryllium and sodium fluorides similar to those
used by Stock, Praetorius et al., whose detailed results had been published the
previous year. The NPL beryllium was 99.5 per cent pure and, like the
German product, was completely brittle when cold.
By 1926 popular interest in beryllium had started to develop. With a
density equivalent to that of magnesium, a stiffness higher than that of steel,
and a melting point approaching 1300°C, this seemed destined to become the
light metal of the future. Towards the end of 1926, when facilities for the
production of relatively pure beryllium existed in Germany, the USA and the
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124

United Kingdom, it was estimated that the metal could, if required, be
produced for approximately 205 per lb. The electrolytic beryllium available in
1927, although reasonably pure, had a Brinell hardness of 140 and was still
brittle when cold. In 1928 when the NPL improved the purity to 99.8 per
cent, beryllium was still brittle and no improvement was noted when the purity
level was gradually increased to the 99.95 per cent level. By 1932 attempts to
produce ductile beryllium at the NPL were abandoned, and it was generally
concluded, throughout Europe and the United States, that beryllium would in
future find its major use as an alloying ingredient where the problems of purity
and ductility were not so critically involved.
Beryllium copper and beryllium nickel alloys were first produced
commercially for Siemens and Halske by Dr W.Rohn at Heraeus
Vacuumschmelze at Hanau, which began to operate just after 1923 at the
height of the inflationary period in Germany. Also in the early 1920s, Michael
Corson in the United States was pursuing similar lines of development. In
1926 he patented the composition of an age hardening beryllium-nickel-copper
alloy, while Masing and Dahl of Siemens also protected the compositions of
beryllium copper and beryllium cobalt copper alloys which had higher
electrical conductivities than those of the American alloy and were therefore of
greater industrial potential.
Beryllium copper displayed a degree of age hardening far higher than that
of any other copper-based alloy, and in spite of its cost beryllium copper soon
found extensive industrial applications. It was particularly valuable as a spring
material: the alloys were soft and ductile in the solution treated conditions and
could then be fabricated into complex shapes. After heat treatment at
temperatures between 300 and 350°C the best alloys developed tensile
strengths approaching 100 tons per square inch, and as a result of the
precipitation which had occurred the electrical conductivities improved to
about 40 per cent that of pure copper. Even higher mechanical properties were
obtained with beryllium nickel alloys, although these were not suited to

electrical work because of their higher resistivity.
DEVELOPMENT OF HIGH TEMPERATURE ALLOYS
One effect of the rapid introduction of mains electricity into the domestic
environment was an increasingly urgent requirement for improved alloys for
electrical heating applications. The alloy required was one which combined a
high electrical resistivity with extreme resistance to oxidation at high
temperatures and mechanical properties high enough to ensure that it did not
fail by creep after prolonged use at a good red heat. Prior to 1900 the only
high resistivity alloys available were cupronickels such as Ferry and
Constantan and iron alloys containing up to 20 per cent nickel.
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Nickel-chromium alloys
The first satisfactory electrical heating alloys, introduced by A.L.Marsh in
1906, were based on the nickel-chromium and nickel-chromium-iron
systems. Wires of these alloys had an electrical resistivity around no
microhms per cm
3
, more than twice that of the best cupro-nickel alloys. They
were, moreover, far more resistant to oxidation and stronger at high
temperatures. Previous investigators had found that chromium additions
tended to increase the oxidation rate of nickel. Marsh found that chromium
additions in excess of 10 per cent rapidly decreased this rate and that alloys
containing around 80 per cent nickel and 20 per cent chromium showed a
good balance between oxidation resistance and resistivity. They were also
ductile enough to be drawn into wire. It is now known that the oxidation
resistance of alloys of this composition depends on the formation of a
protective oxide layer.
These alloys were initially induction melted in air and deoxidized with
manganese. Air melted material was cheap to manufacture, although it was not

always easy to draw into fine wire. At the end of the First World War, Siemens
and Halske found that vacuum melted nickel-chromium alloys were easier to
draw into wire and also had a longer high-temperature working life,
compensating in some degree for the added expense of vacuum melting.
Heraeus constructed at Hanau in 1921 a 300kg (660lb) capacity low frequency
vacuum melting furnace of the Kjellin type for melting nickel-chromium based
resistance alloys. By 1928 two 4000kg (8800lb) furnaces capable of casting two
2000kg (4400lb) nickel-chromium ingots were in regular operation. Siemens
and Halske acquired Heraeus Vacuumschmelze in 1936, after which
considerable quantities of nickel-chromium alloys, beryllium copper, and other
specialized materials for the electrical industry were produced.
The rare earth effect
In the early 1930s it was found that some of the vacuum-melted alloys which
had been processed from melts deoxidized with mischmetall (a rare-earth
mixture) were remarkably resistant to high temperature oxidation. Rohn was
able to show that the metal responsible for this effect was cerium, and in 1934
Heraeus applied for patents covering the manufacture and use of heating
elements to which small quantities of cerium and other metals of the rare earth
group had been added. This ‘rare earth effect’ was followed up in the Mond
Nickel Laboratories at Birmingham, and led to the introduction of the well-
known ‘Brightray C’ series of alloys. The mechanism which permits small
quantities of the reactive metals, such as cerium, yttrium, zirconium,
lanthanum and hafnium to improve the protective nature of the oxides which
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126
form on the surface of nickel chromium and other high temperature alloys is
still imperfectly understood, although the effect is widely employed.
Aluminium containing nickel-chromium alloys
It was found in 1929, by Professor P.Chevenard of the Imphy Steelworks, that
small quantities of aluminium improved the oxidation resistance and high

temperature strength of nickel-chromium alloys and made the alloys responsive
to age-hardening. In 1935 he showed that the strengthening effect of
aluminium could also be augmented by small quantities of titanium.
When, around 1937, Britain and Germany both began to develop a
workable gas turbine for aircraft propulsion, the main technical problem they
encountered was that of producing turbine rotor blades which were strong
enough at high temperatures to withstand the high, centrifugally imposed
tensile stresses. The alloy selected was the 80/20 nickel-chromium solid
solution alloy, strengthened in accordance with Chevenard’s findings by small
quantities of titanium and aluminium. Work on this material started in 1939,
and the first production batch was issued in 1941. Nimonic 80, as the alloy
was called, was the first of the long series of nickel-based ‘superalloys’
produced by Mond. All of these depend for their high temperature strength
upon substantial quantities of the precipitated ? (gamma prime) phase which,
although based on the compound Ni
3
Al, also contains a good deal of titanium.
The first Nimonic alloys were melted and cast in air and hot forged. As the
alloys were strengthened and improved, however, very reactive alloying
constituents were being added, and vacuum melting became mandatory. By 1960,
most of the stronger alloys were being worked by extrusion rather than by forging.
The use of molten glass lubrication, introduced by Sejournet in the 1950s, made it
possible to extrude even the strongest Nimonic alloys down to rod as small as
20mm in diameter in one hot working operation. The alloys were progressively
strengthened, initially by increasing the aluminium content and subsequently by
additions of the more refractory metals such as tungsten and molybdenum.
Nimonic 115, the last of the alloys to be introduced, marked the practical limit of
workability. Stronger and more highly alloyed materials could not be worked and
by 1963 a new generation of nickel-based superalloys was being developed. There
were not worked, but were cast directly to the shapes required.

Since the early years of the century it had been known that the grain
boundaries of metals, although initially stronger than the body of the grains,
tended to behave in a viscous manner as the temperature increased and lost
their strength far more rapidly than single crystals. The concept of an equi-
cohesive temperature, applicable to any alloy, above which the grains
themselves were stronger than the boundaries and below which the boundaries
were stronger than the grains, was first advanced in 1919 by Zay Jeffries. This
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suggested that the strongest high temperature alloys would have large rather
than small grains.
Difficulties were, however, caused on odd occasions by very large randomly
orientated grains. After 1958 it became customary to apply a thin wash of
cobalt-aluminate to the interior of the moulds in which the turbine blades were
cast. This coating helped it to nucleate a uniform grain size in the castings. It
also seemed logical to produce blades having as few grain boundaries as
possible aligned at right angles to the axis of the tensile stress to which the
blade would eventually be subjected. Methods of directionally solidifying
superalloys in such a way that the structures obtained formed a bundle of
longitudinally aligned columnar crystals were first described in 1966 by
B.J.Piercey and F.L. Versnyder of the United Aircraft Corporation. This
process, now known as directional solidification, resulted in an immediate
improvement in high temperature blade performance even without change in
alloy composition. Single crystal blades, which contained no boundaries, were
soon produced as a logical development of the directional solidification
concept. Because grain boundary strengthening additions such as hafnium
were no longer required in single crystal blades, and tended to interfere with
the perfection of their growth, single crystal turbine blades are now
manufactured from alloys which have very much simpler compositions than
the conventional casting alloys.

Cobalt-base high temperature alloys
While the British were using the 80/20 nickel-chromium resistance wire alloy
as a base for their first turbine blade material American manufacturers were
adopting a completely different approach to the problem of high temperature
alloy strength. The blades and other components of gas turbines used for
driving the superchargers of large piston engines had, for a considerable time,
been very effectively and economically manufactured by casting them from
cobalt chromium alloys.
Alloys based on the cobalt chromium system had originally been developed
by Edward Haynes. His original ‘Stellite’ alloy, which would ‘resist the
oxidizing influence of the atmosphere, and take a good cutting edge’ was
patented in 1909. Further improvements in hardness were obtained by
tungsten and/or molybdenum additions. Although originally intended for
dental and surgical instruments, the alloys soon found a considerable industrial
market when it was found that they could be used for heavy turning
operations. During the First World War they were used extensively by the
Allies for shell turning, and in 1918 approximately four tonnes a day of the
alloy were being cast. The Haynes Stellite Company was acquired by Union
Carbide in 1920, and the alloy Vitallium was developed in the late 1920s. This
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128
was successfully cast into turbine blades used to supercharge the engines of the
Boeing Flying Fortress.
The American engine manufacturers understood that no long-term future
existed for workable turbine blade alloys: the high alloying content needed to
achieve high temperature strength inevitably lowered the melting point of such
alloys and simultaneously reduced their workability, and it was illogical that an
alloy specifically designed to withstand creep failure at high temperature
should also be expected to undergo severe plastic deformation during
manufacture. From the very beginning of superalloy development, therefore,

turbine blades were produced in the United States by vacuum melting and
investment casting. The alloys employed were based on the cobalt chromium
system, being strengthened, not by the ‘gamma prime’ phase on which the
British nickel base alloys depended, but by the presence of substantial
quantities of stable carbides.
The way in which the high temperature capabilities of superalloys has
increased with time is illustrated diagrammatically in Figure 1.11. It is
significant that the safe operating temperature for the best single crystal alloys
is only marginally higher than that of directionally solidified material. In view
of the limitations imposed by the known melting points and oxidation
resistance of existing superalloys it seems unlikely that new alloys having high
temperature capabilities greatly superior to those currently available will be
developed. Improved gas turbine performance will, most probably, result from
the exercise of engineering rather than metallurgical ingenuity.
POWDER METALLURGY
Prehistoric iron must have been the first metal to have been consolidated from
a spongelike mass (see Chapter 2). The legend of Wayland the Smith, which
appears to date from the fifth century AD, obviously embodies a good deal of
the folklore on which the whole fabric of powder metallurgy has since been
erected. Various accounts describe him as making a steel sword by
conventional blacksmithing techniques, then reducing it completely to filings,
which were mixed with ground wheat, and fed to domestic geese or hens. The
droppings of these birds were collected, reheated in his forge, consolidated, and
then hammered into a sword. This again was reduced to filings and the
process repeated, the end product being a weapon of unsurpassed strength,
cutting power and ductility.
Apparently even the early smiths were aware that metallurgical quality
could be improved by subdivision, and that by repeating this process of
subdivision a product would be eventually obtained which would combine
hardness and ductility. The improvements which ferrous metallurgists have

attributed to the passage of steel filings through the alimentary canal of a
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domestic fowl are more difficult to account for, however, as it is difficult to
believe that this would reduce the phosphorous content of the metal to any
significant effect. It is possible, however, that the high ammonia content of bird
droppings might help to nitride the metal particles during the early stages of
consolidation, although in the absence of reliable confirmation of this effect it
seems logical to attribute the effectiveness of Wayland’s working approach to
the way in which it was able to progressively refine an initially coarse
metallurgical structure.
The use of powder metallurgy by Wayland the Smith was obviously one
approach to the problem which all swordsmiths have encountered: swords
which were strong and hard had a tendency to break in service, while those
which were free from this defect bent too easily and did not retain their cutting
edge. Wayland’s technique must have produced a refined crystal structure
which, like that which existed in Samurai blades, combined strength and
ductility with the ability to maintain a fine cutting edge. It provides an effective
demonstration of the ability of powder metallurgy to improve the quality of an
existing product even when the metals involved could, if required, be melted
together.
Powder metallurgical techniques were used in pre-Columbian times by the
Indians of Ecuador to consolidate the fine grains of alluvial platinum they were
unable to melt. These platinum grains were mixed with a little gold dust, and
heated with a blow pipe on a charcoal block. The platinum grains were,
therefore, soldered together by a thin film of gold and it was possible by this
method to build up solid blocks of metal which were malleable enough to
withstand hot forging, and were fabricated into items such as nose rings and
other articles of personal adornment.
Although platinum was a metal unknown to the ancients, fragments of this

metal and its congeners (elements belonging to the same group in the periodic
table) were unwittingly incorporated into many of the earliest gold artefacts
from the Old and New Worlds. The first important source of the metal was
Colombia on the north-western corner of South America. Here, the grains of
native platinum were regarded initially as the undesirable component of South
American gold mining operations. The identity of platinum as a new element
was recognized in 1750, by William Brownrigg. In 1754, William Lewis of
Kingston upon Thames found that the grains of platinum could generally be
dissolved in aqua regia, and that the ‘beautiful brilliant red powder’ now
known to be amino-platino-chloride was obtained when sal ammoniac was
added to the aqua regia solution. This precipitate, when calcined, decomposed
to provide a dark metallic powder which appeared to be pure platinum. No
satisfactory method of melting this platinum powder or sponge was found,
however. The first truly ductile platinum was produced in 1773 by Rome
Delisle, who found that if the platinum sponge, after calcination, was carefully
washed and then reheated in a refractory crucible it sintered to form a dull
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130
Figure 1.11
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131
grey mass which could then be consolidated by careful forging at a good red
heat. The density of the mass changed, as a result of this treatment, from 10.05
to 20.12g/cm
3
. This was the first time that the high theoretical density of
platinum had been demonstrated.
This sintering process was described by Macquer in 1778 in his Dictionnaire
de Chimie, and in doing so he provided what is probably the first interpretation
of the scientific principles upon which powder metallurgy depends: ‘a metallic

mass quite compact and dense, but it completely lacks malleability when it has
been exposed to a moderate heat, and only assumes it when subsequently
subjected to a much greater degree of heat’. He goes on to report that:

Particles of platina being infinitely divided in the precipitation it is not surprising
that the heat penetrates such very small molecules more effectively than the
ordinary grains of platina which, in comparison, are of enormous size; and their
softening occurring in proportion, they should show the extraordinary effect on
their agglutination in the proportion of their points of contact; moreover, these
points being infinitely more numerous than can be those of much greater
molecules, solid masses result which have all the appearance of quite dense
metal, melted and solidified on cooling, but they are nothing but the result of a
simple agglutination among an infinite number of infinitely small particles, and
not that of a perfect fusion as with other metals.

This powder metallurgy route was subsequently refined and applied by
William Hyde Wollaston who in 1800 at the age of thirty-four retired from
medical practice and decided to make his fortune by producing and selling
ductile platinum on a commercial scale. Wollaston was a friend of Smithson
Tennant, who between 1803 and 1804 had shown that the black residue which
was left behind after native platinum had been dissolved in aqua regia contained
the two hitherto unknown noble metals, osmium and iridium. The presence of
Figure 1.11: Improvement in the high temperature capabilities of nickel based
superalloys improved since their introduction in 1940. The strongest wrought
alloy to be developed was Nimonic 115 which was introduced in 1960. Stronger
alloys could not be worked, and cast alloys were then introduced. These were
subsequently improved, firstly by directional solidification and finally by turning
them into single crystals
The upper curve on the diagram shows how the solidus of the workable alloys
began to decrease rapidly as the alloys were progressively strengthened by

alloying. By 1960 the gap between the melting point of these alloys and their
working temperature capability had decreased to about 100°C. It was found,
however that single crystal alloys could be based on very simple compositions,
since no alloying additions were required to impart grain boundary strength. As
these alloys developed, therefore, their melting points began to increase, thus
providing a much needed margin of high temperature safety.