miniumof south wales
TRANSCRIPT
University of South Wales"'minium2064812
THE EFFECT OF NEWER METHODS OF PROCESSING ON
THE FATIGUE STRENGTH OF CAST STEEL
I. STRODE, M.Phil., C.Eng., M.I.M. C.G.I.A.
A thesis submitted in pursuance of the requirements of the Council for National Academic Awards, for the degree of Doctor of Philosophy.
Collaborating Establishment. Steel Castings Research and Trade Association, Sheffield.
The Polytechnic of Wales,Department of Mechanical and Production Engineering
FEBRUARY, 1984.
DECLARATION
This dissertation has not been nor is being currently submitted for the award of any other degree or similar qualification.
I. STRODE.
11
ACKNCWLEI)GEMENTS
The author wishes to thank Dr. J.D. Davies, Director, The
Polytechnic of Wales for the provision of laboratory facilities,
also Dr. M.B. Bassett and Dr. C. Davies for their encouragement and
supervision of the work. The helpful discussions held with Dr.
J.D. Griffiths are recorded with gratitude.
The assistance of the technician staff particularly Mr. C.
Monks and Mr. P. Jarman with the machining of specimens is
gratefully acknowledged. The co-operation of Mr. J.C. Thompson,
HIP (Powder Metals) Ltd., Chesterfield with the HIP'ing of
specimens is particularly appreciated.
Finally, my thanks to Mr. M. Jenkins for assistance with
computing, Mrs. C. Tyndell for SEM operation, Mr. A.J. Evans for
graphs and illustrations and to Mrs. H. Hunter for the efficient
typing of the manuscript.
THE EFFECT OF NEWER METHODS OF PROCESSING CN THE FATIGUE STRENGTH OF CAST STEEL
I. Strode, M.Phil., C.Eng., M.I.M., C.G.I.A.
ABSTRACT
The effect of hot isostatic pressing (HIP) and electrochemical machining (ECM) on the microstructure and mechanical properties of cast steel, particularly the fatigue strength, has been investigated.
It is shown that microporosity in cast steel reduces the elongation, reduction in area, and the fatigue strength. However, hot isostatic pressing at an argon pressure of 103MN/m2 and at temperatures varying from 930 °C. to 1210°C. was effective in closing internal microporosity and improving the mechanical properties, particularly the fatigue strength, by up to 70% in a cast low alloy steel. The contribution made by homogenisation of microconstituents to the improvement in the fatigue strength was determined and is shown to be only marginal. The HIP of edge specimens having columnar crystals resulted in an improvement in the fatigue strength. This is attributed to the removal of the anisotropic columnar crystals by isostatic hot working.
The electrochmical machining of wrought and cast steels in a 10% sodium nitrate solution has been carried out. Both the surface finish and the fatigue strength are reduced after ECM and are strongly dependent upon the current density used. It is shown that in spite of their greater heterogeneity and inferior surface finish the reduction in the fatigue strength of cast steels is less than that of wrought steels. The stress relief annealing of mechanically polished specimens resulted in a reduction in fatigue strength of the same order as that obtained by ECM at a high current density. Clearly, ECM produces a surface free from microcracks and compressive stresses. ECM at a low current density similar to that of "stray machining" causes selective attack of the microconstituents with an increased reduction in fatigue strength. However, light shot peening of the surface increased the fatigue strength to a level higher than that of the base metal.
IV
NOMENCLATURE
PS 0.2% Proof stress MN/m2 .
UTS Tensile strength MN/m2 .
EL Elongation%.
R& Reduction in area %.
CVN Charpy V Notch joules.
FL Fatigue limit.
FR Fatigue ratio (FL/UTS).
N Number of stress reversals.
Kf Fatigue strength reduction factor
(FL Plain/FL Notched).
K Elastic stress concentration factor,
q Notch sensitivity factor (Kf~1^(K -1)t
K.. Stress intensity factor Mode 1 MN/m3 / 2 .
K.. Critical stress intensity factor MN/m3 / 2 .
K Maximum value of 1C. during fatigue cycle, nicuc -L
K . Minimum value of K.. during fatigue cycle.
AK Range of K.. during fatigue cycle. (K - K . ) J l ^ ^ max mm
AK Critical value of AK for fatigue crack growth.O
COD Crack opening displacement.
0" Standard deviation
R Surface finish pm a
R Surface finish pm
Xe Electrical conductivity ohm cm
CONTENTS
Page No.
TITLE.
DECLARATION.
ACKNOWLEDGEMENTS.
ABSTRACT.
NOMENCLATURE.
CONTENTS.
CHAPTER I
CHAPTER II
CHAPTER III
CHAPTER TV
CHAPTER V
CHAPTER VI
CHAPTER VII
INTRODUCTION.
LITERATURE SURVEY.
EXPERIMENTAL PROCEDURE.
RESULTS - HOT ISOSTATIC PRESSING.
RESULTS - ELECTROCHEMICAL MACHINING.
DISCUSSION - HOT ISOSTATIC PRESSING.
DISCUSSION - EI.ECTROCHEMICAL MACHINING.
CHAPTER VIII CONCLUSIONS AND FURTHER WORK.
REFERENCES.
REPRESENTATIVE COMPUTER GRAPHS.
WEIBULL DISTRIBUTION.
APPENDIX I
APPENDIX II
APPENDIX III
APPENDIX IV
APPENDIX V
APPENDIX VI
APPENDIX VII
GRAPHS AND PHOTOGRAPHS RELATING TO CHAPTER III.
GRAPHS AND PHOTOGRAPHS RELATING TO CHAPTER IV.
GRAPHS AND PHOTOGRAPHS RELATING TO CHAPTER V.
GRAPHS AND PHOTOGRAPHS RELATING TO CHAPTER VI.
PUBLICATIONS.
Al.
A2.
A3.
A4.
A5.
A6,
l
ii
iii
iv
v
vi
1
3
36
43
54
59
72
79
83
,1 - A1.10
,1 - A2.3
,1 - A3.7
,1 - A4.32
,1 - A5.21
,1 - A6.3
vi
CHAPTER I
Introduction
In recent years, increasing demands have been made for steel
castings with a greater freedom from defects for use under conditions of
high stress.
Much progress has been made to achieve this objective, by the use
of foundry techniques designed to produce directional solidification
toward the feeder heads. The success of these methods is evidenced by
the increasing use of cast steel components for important engineering
applications.
2 3 It has been shown by a number of investigators, ' that the
mechanical properties of cast steel generally decrease with increasing
section thickness, particularly the elongation and reduction of area
values. To minimise this problem, methods of unidirectional
solidification have been developed, which generally use water cooled
metallic sections for critical parts of the casting. These have been
particularly successful in the case of investment cast superalloys for
4 5 gas turbine components. ' However, the wider application of these
methods is restricted by design limitations, casting methods, and high
cost.
An alternative method of combating internal porosity in cast metals
is now possible, by subjecting solid castings to the process of hot
isostatic pressing. A uniform stress is applied by means of high purity
- 1 -
argon gas which induces the collapse of internal cavities. The
operation is carried out at an elevated temperature which completes the
closure of the pores by diffusion bonding.
It is being increasingly appreciated that the mechanical properties
of cast metals may also be impaired by the final machining operations
which are carried out. The concept of "surface integrity" is now being
applied, which includes the examination of the surface structure,
surface residual stresses and surface properties, as well as the
conventional surface finish measurements.
A number of sophisticated methods of metal removal are now in
g general use, such as electro-discharge machining (EDM) and electro-Q
chemical machining (BCM), in addition to the more traditional methods
of metal cutting and grinding. The effect of these newer methods of
machining on the surface integrity of cast metals has been little
investigated.
In the present work the use of hot isostatic pressing (HIP) to
eliminate internal microshrinkage cavities and the effect of electro
chemical machining on the surface properties of cast steel has been
investigated.
Since the fatigue strength of steel is particularly structure
sensitive, special emphasis has been given to the effect of these
processes on the fatigue strength.
- 2 -
CHAPTER II
LITERATURE SURVEY
2.0 HOT ISOSTATIC PRESSING
2.1 Introduction
The process was developed at the Battelle Columbus
Laboratories in 1955, originally as an isostatic diffusion bonding
process for cladding nuclear fuel elements. It was later
extended to the itianufacture of powdered components particularly of
the more difficult metals to fabricate such as beryllium, titanium
and the nickel and cobalt base superalloys.
The potential use of hot isostatic pressing for the closure of
internal porosity in cast metals was also realised, and has been
increasingly applied to both ferrous and non-ferrous alloys for
this purpose.
2.2.0 NON-FERROUS ALLOYS
2.2.1 COPPER BASE ALLOYS
Bronze castings have a long freezing range and are particu
larly prone to interdendritic porosity especially when cast in
thick sections. However, hot isostatic pressing has been success
fully applied to bronze castings for nuclear submarine fluid
transfer systems. After HIP treatment at temperatures ranging
- 3 -
from 677 to 815°C at an argon pressure of 103MN/m2 for three hours,
increases in yield stress of 14%, U.T.S. of 37%, and elongation of
100% were obtained compared with the "as cast" properties. In
addition to the effective closure of porosity, interdendritic
segregation was reduced and second phase particles were
redissolved.
2.2.2 ALUMINIUM AUDYS
Considerable interest has been shown in the application of
hot isostatic pressing to aluminium alloys for use in the
automotive industry.
Initial investigations by the Aluminium Company of America
resulted in the development of the Alcoa 359 process which was
12 subsequently patented. Both independent test bars and commercial
castings in alloys A356-T6, A357-T6 and F132-T6 were subjected to
HIP at a pressure of 103MN/m2 for 2-3 hours at a temperature near
to the solution temperature of the alloy. As a result, internal
gas and shrinkage cavities were completely closed with a resulting
increase in the yield stress, U.T.S., and % elongation, depending
in magnitude on the alloy composition. The most significant
improvement was an increase of up to 300% in the fatigue strength
compared with that of untreated castings, and approaching that
obtained in comparable forged products. Further investigations,
showed that the presence of hydrogen gas in the castings tended to
inhibit the effectiveness of the HIP treatment in closing porosity.
This emphasised the importance of the effective degassing of
aluminium castings prior to HIP.
- 4 -
The benefits to be obtained from HIP have been shown in the
development of a new cast aluminium alloy A201-T7, with mechanical
14 properties comparable to wrought alloys. The widespread
application of this alloy was inhibited by the presence of porosity
causing a wide scatter in the mechanical property values. After
HIP, the microshrinkage porosity was completely eliminated with
consequent improvement in mechanical properties, and reduction in
data scatter. A marked improvement occurred in the % elongation
particularly in thicker sections; also the fracture toughness
values increased from 35 to 44MN/m 2 . The notch fatigue (Kt 3.0)
and the fatigue crack propagation rate was slightly superior to
that of the contending wrought aluminium alloys. The effect of HIP
on alloy A201 has been further evaluated. ' Both U.T.S., and
high cycle fatigue tests showed a marked improvement with closure
of internal porosity. Radiographic grade C castings showed almost a
two fold increase in fatigue limit after HIP, upgrading to radio-
graphic grade A, accompanied by a significant reduction in data
scatter.
An important property in the evaluation of aluminium alloys
for automotive cylinder heads is thermal fatigue. The presence of
porosity causes a reduction in ductility which is detrimental to
thermal fatigue. It has been shown that HIP was effective in
reducing microshrinkage cavities in aluminium alloy (G-AlSi9Cul).
This resulted in increased ductility values and a 50% improvement
in thermal fatigue resistance, together with a change in fracture
mode from brittle to ductile fracture.
The HIP of aluminium alloys is conducted in the region of the
solution temperature and a summary of the reported HIP parameters
is given in table 2.1- 5 -
TABLE 2.1
REFERENCE NO.
13
13
181915
ALLOY TYPE
A356-T61 A357-T62F132-T6 F142-TF
-A357A201-T7
TEMP. °C.
-
-
510510510
PRESSURE MN/m2
103
103
10370
103
TIME hrs.
2-3
2-3
226
2.2.3 MAGNESIUM AND THORIUM
Whilst no detailed results are available, it is stated that
these metals may be HIP'ed under the same conditions as aluminium
alloys namely at 510°C/70MN/m2 /2 hours. 19
2.2.4 TITANIUM ALLOYS
The application of HIP to titanium alloys has been particularly
successful since "the difficulty of super heating this reactive
metal makes complex castings almost impossible to feed foror\
soundness". By means of HIP, the ductility and fatigue properties
are markedly improved without loss of yield strength or U.T.S. In
one foundry with HIP facilities about 95% of titanium alloy and
other castings are now being "HIP'ed".
The HIP of Ti-6Al-4V alloy castings was first conducted at
Battelle. 10 Full densification was obtained by HIP at
968°C/69MN/m2 /l hour. Incomplete void closure was obtained with a
- 6 -
reduced time of 0.5 hours or with a reduced temperature of 871°C/3
hours. Using the original parameters the high cycle fatigue
strength at a temperature of 316°C was considerably improved as
well as the stress rupture life at 400°C. Since surface connected
porosity remains unaffected after HIP, a number of sealing
techniques were tried. Electron beam welding, TIG welding and
vacuum encapsulation were found to be effective.
The influence of HIP temperature on cast Ti-6Al-4V alloy has
21 been further investigated. For a constant pressure of 103MN/m2
and time (1 hour), a reduction in temperature from 954°C to 843°C
resulted in increased strength, but gave an unpredictable decrease
in ductility in some cases. This was attributed to insufficient
bonding of the collapsed pores.
A number of investigators have reported that whilst HIP has
little effect on the yield stress and U.T.S., both the low cycle
fatigue and high cycle fatigue of titanium alloys are improved. An
22 example of a Ti-6Al-2Sn-4 Zr-2Mo alloy is given by Widmer.
The fatigue strength of Ti-6Al-4V alloys at 20 °C is increased
19 by 25% after HIP plus an unspecified heat-treatment. In addition
to an increase in the statistical mean value of the fatigue life, a
reduction in data scatter by a factor of six is reported. This
reflects the increased reliability of the product after HIP.
The influence of the subsequent heat treatment of HIP'ed
23 Ti-6Al-4V cast alloys is shown by Freeman. A marked increase in
the fatigue strength at 20°C after HIP plus annealing was obtained.
- 7 -
However, a substantial increase occurred when the castings were
solution treated and aged after HIP. Again the reduction in data
scatter is evident.
A more detailed study of the effect of HIP on fatigue strength
24 has been conducted by Teifke et.al. It is shown that the as cast
fatigue limit at 5 x 10 cycles increased from 275MN/m2 to 415MN/m2
after HIP. However, further solution heat treatment and ageing
resulted in lower fatigue values. The microstructure revealed alpha
plates at the prior grain boundaries even after solution treatment
at 1005°C. These had previously been shown to be favoured sites
for the initiation of fatigue cracks.
o/rIn further work Elyon found only a slight improvement in the high
cycle fatigue of Ti-6Al-4V castings after HIP and subsequent
annealing. This was not due to insufficient porosity closure, but
to crack initiation at large alpha colonies at grain boundaries,
particularly in the previously existing porous areas that had been
healed by HIP. In view of the conflicting evidence produced by
23 Freeman and others it is clear that further research is required,
particularly the effect of post HIP heat treatments.
Due to the increasing use of titanium alloys for gas turbine
27 components, the mechanical properties at elevated temperatures are
important.
The work at Battelle showed that the high cycle fatigue
strength at a temperature of 316°C was considerably improved after
18 HIP. Further, Bailey and Schweikert found that the fatigue limit
- 8 -
(10 cycles) at 871°C increased from 310MN/m2 in the as cast state
to 414MSI/m2 after HIP. Also, the low cycle fatigue at 427°C after
HIP was similar to that of forged material. All specimens were
given the same final heat treatment.
By the use of HIP, a cast-to-net-shape martensitic transage
titanium alloy has been developed as a replacement for a wrought
28 Ti-6Al-4V alloy for aircraft engine rotating components. Transage
134 and 175 cast-to-size test specimens were subjected to HIP and
subsequently aged at 538°C. These alloys showed a superior yield
stress and U.T.S. than wrought Ti-6Al-4V components at temperatures
up to 500°C. However, the yield stress of the HIP specimens was
directly related to the cooling rate from the solution heat-
treatment temperature. The transage 175 alloy showed a distinct
4 5 fatigue limit at between 10 and 10 cycles. The axial fatigue
stress corresponding to 10 cycles was comparable to wrought
Ti-6Al-4V sheet specimens at 121°C, and was superior at 260°C.
The elevated temperature creep properties are also important
and the stress to rupture values of cast plus HIP Ti-6Al-4V alloys
are superior to that of the as-cast alloys after a similar heat
treatment. ' Parity with comparable wrought alloys is obtained.
A comprehensive investigation into the creep behaviour of Ti-6Al-4V
alloys in the as-cast condition, and after HIP, within the99
temperature range of 121°C to 454°C, has been reported. For creep
tests up to 316°C only primary creep occurs. The creep rates
decrease with time, and after about 200 hours become equal to zero
which is known as the "creep saturation" condition. In the case of
cast plus HIP tests the total plastic strain at which creep
- 9 -
saturation takes place increases steadily with increase in test
temperature. Also the time taken to reach creep saturation
decreases as the temperature increases from 121°C to 177°C, after
which it remains within 100 to 200 hours. The effect of grain size
is important. For a given temperature and stress level, the
saturation plastic strain is greater and the time taken in reaching
saturation creep less for a fine grain size when compared with
specimens having a coarse grain.
Some investigators have reported a significant scatter in the
mechanical properties of both as-cast and cast plus HIP Titanium
alloys. After the HIP of Ti-6-2-4-2-S thin gauge investment
castings, the U.T.S. varied from 926.8MN/m2 to 1082.5MN/m2 and the %
elongation on a 25.4 mm gauge length from 4.0 - 11.0%. The data
scatter was not due to residual porosity since microscopic
examinations confirmed its absence. Scanning Auger microscopy
revealed considerable alloy segregation particularly at the sites of
healed pores which persisted after 10 hours annealing. The data
scatter may be due to such alloy segregation which may be indigenous
to this alloy composition. In the case of cast and HIP Ti-6Al-4V
31 alloys, Larson and Wright have reported that a minimum U.T.S. of
900MN/m2 , 6% elongation (min) and 10% reduction in area is
consistently achieved. Other investigators have also reported
21 23 increases in % elongation after HIP. '
The HIP of Titanium alloys is carried out below the 995 °C
transition temperature, and the reported parameters are summarised
in table 2.2.
- 10 -
TABLE 2.2
REFERENCE NO.
10
18
18
19
21
23
24
26
28
29
30
31
ALLOY
T1-6A1-4V
T1-6A1-4V
Ti-6-2-4-2
T1-6A1-4V
T1-6A1-4V
T1-6A1-4V
T1-6A1-4V
T1-6A1-4V
Transage
134 and 175
T1-6A1-4V
Ti-6-2-4-8
T1-6A1-4V
TEMP. 0 C
968
900
900
968
954
815-955
900
900
815
900
900
900
PRESSURE MN/m2
70
103
103
70
103
103
103
105
103
103
-
103
TIME hrs.
1
2
2
1
1
2-7
4
2
2
2
2
-
2.2.5 SUPERALLOYS
HIP has been extensively applied to superalloys for gas turbines
and aeroengines where porosity in highly stressed components may
result in premature failure due to fatigue and/or creep.
The first comprehensive work was reported by Wasielowsky and
Lindblad. 32 The HIP pressure was varied from 35MN/m2 to 207 MN/m2
but had little effect above a critical value (presumably the
conpressive yield stress). Twinned crystals in the post HIP micro-
- 11 -
structure was indicative of plastic deformation during void closure.
The temperature required for complete densification varied with
alloy composition. In the case of Rene 77 and IN-738, a HIP
temperature of 1098°C to 1177°C produced only 50 to 80% void closure,
and a temperature of 1177°C to 1204°C was necessary for 100%
densification. A higher temperature was required for the stronger
alloy IN-792.
During HIP, some of the gamma prime (y 1 ) was agglomerated due to
slow cooling from the HIP temperature. This impaired the mechanical
properties which necessitated a four stage post HIP heat-treatment in
order to restore the normal morphology of the Y phase.
After the post HIP heat treatment, the ambient temperature YS
and UTS of IN-738 was little affected, but the tensile ductility
increased by a factor of 3. At elevated temperatures the stress
rupture life and ductility was increased. The Larson-Miller
parameter showed a 50% improvement in data scatter of the average
rupture life at 982°C/152MN/m2 and over 250% improvement at
871°C/276MN/m2 . The low cycle fatigue life at 816 °C was increased
by a factor of 8, but the high cycle fatigue life was little
affected.
Van Drunen et.al have shown that by controlling the rate of
cooling from the HIP temperature, the four stage heat treatment may
be replaced by a single low temperature ageing treatment, thus
improving the cost effectiveness of the HIP process. At
760°C/586MN/m2 both heat treatments gave similar stress rupture
- 12 -
values, but at 830°C/345MN/m2 the modified heat treatment gave
superior values. At higher temperatures the difference between the
heat treatments tended to diminish. It is therefore evident that
34 alloy IN-738 is subject to the "ductility trough" problem, but is
less affected by the modified heat treatment.
Actual turbine blades subjected to HIP and the modified heat-
treatment showed improved 0.2% YS, UTS and %RA at 20°C and 650°C, but
the % elongation was little affected. Also the stress rupture life
was increased by* 30% with a reduced Larson-Miller scatter band.
30 In contrast to that previously reported, the high cycle
fatigue of IN-738 was increased by between 10% to 20% after the
modified heat treatment. The improved properties were due to the
production of serrated grain boundaries which are normally
characteristic of wrought superalloys. An improvement in the high
cycle fatigue at 850°C of Nimocast alloy 738LC after HIP, plus a
standard heat treatment has also been reported by McColvin .
However, little change in the 0.2% PS, UTS, EL% and RA% at 20°C and
850°C or stress rupture values at 850°C occurred.
The addition of hafnium to nickel base superalloys improves
ductility and oxidation resistance at high temperatures, but the
"castability" is impaired. The greater susceptibility of those
alloys to porosity makes them particularly suitable to improvement by
HIP.
23 37 It has been reported ' that IN-792 containing hafnium,
HIP'ed at 1150°C to 1230°C/103MN/m2 /2-7 hours, showed a marked
- 13 -
increase in the UTS and tensile elongation at 20°C, but little change
in the yield stress occurred in specimens machined from 127 ram thick
sections. The elevated temperature properties were also improved
with an almost forty-fold improvement in the -20 rupture life at
760°C/586MN/m2 and 871°C/345MN/m2 . Also a significant increase in
the low cycle fatigue at 480°C occurred. For lives over 104 cycles,
the data agreed with the analytical predictions of the Manson-Cof f in
equation which was originally derived for wrought alloys.
23 The effect of section size has been shown by Freeman , since
the marked improvement in the stress rupture values at 980°C/200MN/m2
after HIP was not sustained in thicker sections of alloy B-1900+Hf,
and were reduced in the IN-792+Hf alloy. The effect of HIP on actual
production components is therefore a more realistic approach than the
use of small independent test bars.
The mechanical properties of sectioned IN-713C turbine castingsop
have been reported. After HIP at 1232°C/107MN/m2 /4 hours and
subsequent heat treatment, the tensile ductility and fracture
toughness at 20°C was slightly improved, but the low cycle fatigue
life at 649°C was substantially increased. However, contrary to
expectation, the high temperature stress rupture 982°C/152MN/m2 was
39 virtually unaffected. Lamberigts et.al, have shown that a rapid
rate of cooling from the HIP temperature is essential. After the HIP
of IN713C turbine blades and cooling at 1000°C/hour between 1220°C
and 650 °C the creep life at 760°C/530MN/m2 was superior to the
normally treated blades. A satisfactory microstructure was produced
after HIP, which made a post HIP ageing treatment (930°C/16hours/air
cool) unnecessary and possibly damaging to the production of optimum
mechanical properties.
- 14 -
Cast superalloy Rene 77 and Rene 80 can be fully densified by
HIP at 1200°C/69MN/m2 /4 hours or at 1218°C/103MN/m2 /2 hours. 18 ' 32 ' 40
Significant increases in stress rupture of Rene 77 at 815°C/138MN/m2
and 980°C/152MN/m2 and Rene 80 at 870°C/310MN/m2 was reported. The
high cycle fatigue (HCF) at 107 cycles of Rene 80 at 871°C increased
from 320MN/m2 to 400MN/m2 after HIP. 41 The densification of Rene 120
was incomplete after HIP at 1177°C/103/MN/m2/4 hours, but porosity
was reduced to zero after the temperature was increased to
22 42 1204°C. ' Whilst the UTS was only marginally improved, the YS
increased from 580 to 610MN/m2 , elongation from 2.5 to 4.2%, and the
average cycles to failure from 1,850 to 12,080 at a stress of
586MN/m2 at 871 °C. As a result of these investigations, an
integrated casting plus HIP process was adopted for the manufacture
of cast superalloy gas turbine components.
The importance of establishing optimum HIP parameters for each
alloy composition has also been shown for the Inconel range of
18 alloys. A standardised HIP pressure of 103MN/m2 /3 hours was used
for all alloys. A temperature of 1163°C was sufficient to completely
densify Inconel 718, but 1190°C was required for Inconel 738 and
Inconel W. Whilst there was no increase in the 0.2% YS and UTS at
20°C of IN-718, the reduction in data scatter was considerable. '
However, the % elongation and % RA was substantially increased. With
a modified post HIP heat treatment, a substantial increase in the OTS
of IN-718 was obtained both at 20 °C and at increasing temperatures up
1 Q to 650°C. The level of UTS approached that of the forged alloy.
Schweikert has also reported a substantial increase in the 0.2% YS
and UTS of IN718 castings after a revised post HIP heat treatment.
The statistical scatter was reduced, the -30" limit was increased by
- 15 -
a factor of 50, resulting in a 50% increase in the design allowance
45 for HIP'ed castings for gas turbines.
The microstructural changes which occur during the HIP of cast
Inconel 718 has been investigated by Bouse and Shilke. In addition
to porosity closure, homogenisation also occurs, and the niobium rich
Laves phase diffuses into the matrix, so freeing the niobium for the
y phase. As a result the 0.2% YS and UTS increased by 10 - 20% up
to a temperature of 650°C, and the %RA up to 450°C. However, the %
elongation was reduced by up to 28%.
Special attention to the HIP cycle is necessary in the case of
superalloys containing refractory elements molybdenum and tungsten
such as M&R-M-246. Erratic results have been obtained due to the
38 formation of M,C and sigma phase. In order to obtain improved
creep rupture properties at 760°C/672MN/m2 , the post HIP heat
treatment was preceded by heating for two hours at 1218°C, which was
higher than the HIP temperature, followed by air cooling. The
addition of MC stabilisers such as hafnium or niobium was recommended
47 if HIP was adopted as an essential part. However Burt et.al have
shown that when the rate of cooling from the HIP temperature of
MAR-M-002 (10%W) was controlled to produce a y' distribution
comparable to that of the cast alloy, improved creep ductility was
achieved without adversely affecting the creep and fracture
39 resistance at 950°C/130-250MN/m2 . Similarly, Lamberigts et.al,
achieved improved stress rupture values at 760°C/692MN/m2 in
MAR-M-002 with optimum HIP parameters and a controlled rate of
4fi cooling. Viatour et.al have also reported increased stress rupture
and creep ductility at 760°C/695MN/m2 of a MAR-M-002 (10%W, 1.5%Hf).
- 16 -
The HIP cycle was followed by a single ageing treatment 870°C/16
hr/air cool, and the improved properties were related to the -y'
morphology and distribution which was dependent upon the cooling
rate below 1200°C.
HIP has also been used to improve the mechanical properties of
cobalt base alloys for surgical implants. After HIP at
1230°C/103MN/m2/2 hours followed by solution treatment at 1250°C plus
air cooling, a Haynes Stellite No. 21 alloy showed no change in YS
and UTS but the % elongation and % RA increased from 5 - 17%. The
fatigue strength and crevice corrosion susceptibility was also
49 improved. However, a similar alloy after a further ageing
treatment at 650°C/2 hours showed an increase in YS and UTS as well
as the % elongation. Stellite extrusion dies have also been HIP'ed
successfully. When backed with high speed steel, these dies have a
good heat resistance with an increased die life of from 6-8 times
52 that of standard tool steel dies.
To obtain efficient pore closure, the HIP temperature for super-
alloys should be between the solvus and incipient melting. A summary
of the reported HIP parameters for superalloys is given in table 2.3.
- 17 -
TABLE 2.3
REFERENCE No.
183233
36
381819
46
183223
18,32,4118,32,41
1919
22, 423847483938
49
50
ALLOY TYPE
IN-738IN-738IN-738
NIMOCAST - 738LCIN-713CIN-718IN-718
CAST ALLOY 718
INCONEL WIN-792
IN-792HfRENE 77, 80RENE 77, 80RENE 80RENE 120RENE 120MAR-M-002MAR-M-002MAR-M-002MAR-]YK>04MAR-M-246
HAYNES-STKT .T .ITE NO. 21
HAYNES-STELLITE NO. 21
TEMPERATURE°c.
11901177 - 1205
1200
1200
123211631177
1172 - 1213 (1200)1190
1205+1150 - 1230
12001218120012181205117012301205
1170-12201205
1230
1200
PRESSURE MN/m2
103-100
103
10710370
103
103-
10370
10370
103103140103103140107
103
102
TIME HRS.
31 (min.)
2
4
433
2-4
31 (min.)2-7
42444444442
4
HIP has also been used to "rejuvenate" superalloy turbine
components which have reached their calculated creep rupture life.
Internal creep voids are closed resulting in the restoration of
stress-rupture values, and low cycle fatigue life, to almost their
original level. Preliminary welding, EDM machining, and surface
coating may be necessary as well as post HIP heat treatment. HIP
has been applied in this way to nickel base superalloys IN-718,
INCOLOY 901,54 MAR-M2 (directionally solidified), WASPALOY,
Ti-6Al-4V turbine disks, and bronze components. However, in
some cases, regeneration of high temperature creep properties may be
achieved by periodic heat treatment only.
- 18 -
2.3.0 FERROUS ALLOYS
Compared with superalloys, less attention has been given to the
HIP of ferrous castings apart from some precipitation hardened (PH)
stainless steels.
2.3.1 STAINLESS STEELS
Complex stainless steel castings are susceptible to internal
shrinkage porosity resulting in high rectification costs. HIP is
therefore being used as a means of improving the mechanical
properties of cast CF-8 and CF-8M castings for nuclear reactorCO
components. The YS, UTS and % elongation were improved together
with up to 95% saving in rectification costs. Where surface
porosity occurs an effective sealant is required prior to HIP.
P^amjet inlet 17-4PH castings have been HIP'ed at
1120°C/103MN/m2 /2 - 4 hours, followed by solution treatment at
59 1038°C and ageing at 538°C. No change in the YS and UTS occurred
but the elongation % and RA% was increased by 31% and 27%
respectively. The fracture toughness in air (Kj ) and in a hostile
environment (1C. ) after HIP was increased to the same level as the Tec
wrought alloy. The improvement in mechanical properties after HIP
was attributed to the dispersion of ferrite "stringers".
The high cycle fatigue (HCF) of 15-5PH investment castings
after HIP and solution treatment at 1040 °C plus ageing at 540 °C, was
increased by 91% and 25% when compared with specimens cast in a
production shell and a high conductivity shell respectively. The
HCF of notched (K 3.0) cast-to-size specimens was increased by 100%
- 19 -
after HIP, and specimens machined from helicopter lag damper
castings by 140%. Microscopic examination revealed the absence of
both porosity and delta ferrite after HIP .
HIP has also been successfully applied to the improvement of TIG
welds in AISI type 304 stainless steel using a type 308 filler
metal. 61 ' 62 After HIP at 1040 - 1095°C/103MN/m2/l - 3 hours,
internal microporosity and other welding defects were eliminated.
Whilst the UTS was little affected, the elongation was increased by
140% in pipe welds and 173% in plate welds at 20°C. However, a
reduction of about 40% in YS occurred, but the final value was within
the minimum ASME requirements. The Charpy V notch values of the type
304 base metal was increased by 70% after HIP, 147% in the HAZ and
235% in the fusion zone. In addition to void closure, HIP resulted in
homogenisation of the fusion zone, a reduced number and size of
"stringer" inclusions and a reduction in the ferrite number (FN) from
12.3 to 1.7. The average FN for normal welds was 11.1% and the
reduction after HIP accounts for the reduced YS.
HIP has proved beneficial when applied to hard facing alloys such
as Tribaloy 800 on type AISI 316 stainless steel. Greater wear
resistance and reliability has been obtained in the plugs and seats
used in a liquid sodium environment at temperatures up to 650°C in
nuclear breeder reactors.
2.3.2 OTHER FERROUS ALLOYS
HIP has also been applied to an 18% Nickel Maraging steel. An
increase in YS, % elongation, and % RA was obtained.
- 20 -
Mainly for economic reasons, the application of HIP has been
largely confined to highly alloyed material. However, a growing
,63interest is reported in its use for the up-grading of low alloy
steels which has been investigated in the present work.
Certain grades of cast iron benefit from HIP. The UTS of a
high chromium iron casting was increased from 627MN/m2 in the
as-cast state to 1040MN/m2 after HIP. 64
The reported HIP parameters for ferrous castings are given in
table 2.4
TABLE 2.4
EEFERENCE NO.
1819596061
62
23
ALLOY TYPE
17-4PH17-4PH17-4PH15-5PH
MSI 304/308TRIBALOY 800 / AISI 41618Ni MARAGING
TEMPERATURE °C
1066117711201120
1040 - 1095
1093
1065 - 1205
PRESSURE MN/m2
10370
103103103
103
103
TIME HRS.
2343
1-3
12-7
2.4.0 ECONOMIC FACTORS
When HIP is used as an integral part of the processing route,
benefits other than improvement in mechanical properties accrue, which
reduce the overall cost of the process.
2.4.1 FOUNDRY PRACTICE
Since internal microporosity will be subsequently closed by HIP,
casting design can be simplified and gat ing/feeding systems modified.
- 21 -
A lower casting temperature may be possible which coupled with a
reduced charge weight, will result in a saving in fuel costs.
Acceptable properties can only be obtained in some superalloys by
the use of virgin charge materials. However, with the inclusion of
HIP, improved properties are obtained with the use of lower cost
39 revert material. Further economies are obtained from a reduction
in the degree of non-destructive testing required, a lower rejection
rate, and reduced rectification. The rejection rate of a highly
stressed stainless steel impeller was reduced from over 90% to
20 almost zero after HIP. Similarly for Rene 120 turbine blades, a
scrap rate of 28% was reduced to 4% after HIP. 19
2.4.2 NEAR-NET-SHAPE PHILOSOPHY
Advantage may be taken of the greater dimensional accuracy and
reduced machining allowances of investment castings wherever
possible. Components which were previously forged or fabricated may
be cast as complete components with considerable cost savings. A
helicopter main rotor hub casting (Ti-6Al-4V) which was HlP'ed,
41 replaced a forging with a saving of about 20% in metal usage.
2.4.3 IMPBOVED WELDABILITY
Some alloys which in the as-cast state were prone to weld
cracking, e.g., IN-718, have shown improved weldability after HIP.
An engine mount previously made as a weldment from wrought
components is now made as a series of IN-718 HIP'ed castings which
are subsequently welded together. Minimal HAZ weld cracking
44 66 occurred and no post heat treatment weld cracking. '
- 22 -
2.4.4 IMPROVED MACHINABILITY
Superalloys in particular are difficult to machine by conven
tional methods, but IN-718 after HIP showed improved
44 machinability. The use of chemical milling as a finishing process
was impracticable in the case of many cast superalloys due to
microporosity and microstructural inhomogeneity. However, after
HIP, several alloys such as IN-718, Ti-6Al-4V and 17-4PH stainless
steel castings could be chemically milled at uniform rates resulting
in acceptable surface properties. '
2.4.5 HEAT-TREATMENT COSTS
Many superalloys require a multi-stage heat treatment after HIP
in order to produce the required microstructure and mechanical
properties. This is mainly due to a slow rate of cooling from the
HIP temperature. However, improved mechanical properties can be
produced by a modified heat treatment which is incorporated into the
HIP cycle followed by a single ageing treatment. In some cases,
no post HIP heat treatment is required provided that the rate of39 cooling is sufficiently fast to simulate the casting conditions. ,
One of the projected developments designed to increase the cost
effectiveness of HIP is the building of larger units equipped with
heat exchangers that allow part cooling under constant pressure. In
addition to a 50% reduction in cycle time, direct quenching will be
20, 67 possible from the solution treating temperature.
- 23 -
2.4.6 MATERIAL DEVELOPMENT
New alloys are being developed with HIP as an integral part of
14 15 the production route. Aluminium alloy A201 ' and the transage28 Titanium alloy are typical examples. In the manufacture of many
cast components, HIP is now an essential part of the process.
2.5.0 ELECTROCHEMICAL MACHINING
2.5.1 INTRODUCTION
ECM is a method of metal removal by initiating electrolytic
action. Using a suitable electrolyte a low voltage D.C. current is
passed between the component which forms the anode and a purposely
shaped cathode, which results in the controlled dissolution of the
anode. The gap between the anode and cathode is about 0.125 mm
through which the electrolyte is passed at velocities as high as
3000 to 6000 on/sec. to remove the reaction products.
The rate of metal removal is determined by Faraday's Law and
is proportional to the local current density which for most
metals is about 8-16 cmVmin., for a current of 10,000 amperes.
To obtain high metal removal rates, current densities of up to 200
A/cm2 or more are used. 68 The local rate of metal removal is
highest at the points of closest approach of the electrodes, and as
the tool is advanced by a precise mechanism toward the workpiece,
the shape of the workpiece conforms closely to that of the tool.
- 24 -
The ECM process was developed mainly at Battelle in the U.S.A. 69
and at PEPA in the U.K. The technique is now well established as a
production method with well defined limits of application. 71 ' 72 The
purpose of the present review therefore is to consider the effect of
ECM on the surface integrity of ferrous materials. The main emphasis
seems to have been directed at producing an acceptable surface finish,
but the effect of ECM on the surface microstructure and mechanical
properties is also important.
2.5.2 SURFACE INTEGRITY
When compared with other methods of metal removal, properly
conducted ECM is capable of producing a good surface finish. An Ba
of 0.10 ym to O.SOum may be consistently produced but both the surface
finish and surface integrity is affected by the current density, anode
potential, type, concentration, temperature and characteristics of the
electrolyte as well as the workpiece material. These factors are
often interrelated.
2.5.3 CURRENT DENSITY AND ANODE POTENTIAL
These are major factors in controlling the surface structure of
the workpiece and much information may be gained by means of
laboratory determined Potentiostatic curves. The surface finish will
vary depending on whether the ECM is carried out in the etching,
polishing, passive or transpassive part of the anode potential/current
density curve. 7 Provided that other factors remain constant, the
surface finish generally improves as the current density increases.
This has been shown for a Nimonic 80 electro-machined in a saturated
- 25 -
solution of sodium chloride at 20°C, and for a 0.45% carbon
7fisteel/sodium chloride solution.
A low current density may result in etching the surface to a
depth of approximately 0.015 mm, with a consequent deterioration in
surface finish. Since current density decreases rapidly with
distance from the cathode face, the surface finish of component side
walls is inferior to that of the frontal surfaces of the cathode. The
side wall surface may be etched or have a matt appearance with a
78 surface finish of up to Sum Ra. At low current densities "pitting"
of the surface may occur resulting in an inferior surface finish.
Some materials form passive films which are attacked by aggressive
anions such as chlorides. This leads to localised break up of the
79 film and concentrated attack forming a "pitted" surface. Pitting
corrosion may also occur at a higher current density during the
80 "polishing" stage due to the onset of gas evolution.
In alloys that passivate, an increase in current density alters
the anode characteristics from a passive to a transpassive state and
efficient ECM takes place resulting in a good surface finish. For
example, the ECM of a low alloy steel in a sodium chlorate solution at
current densities of 39 to 116A/cm2 and a maximum flow rate of 0.63 x
81 10 *m3 /s produced a surface finish of 0.10 ym. However, a sudden
change from the passive to the transpassive state results in an
82 inferior surface finish due to pitting. The surface finish (Ft)
changed from Sum in the active region to 50pm in the transition zone
before stabilising at 1pm in the transpassive condition.
- 26 -
2.5.4 ELECTROLYTES
The surface finish of the workpiece is also dependent upon the
type, concentration and temperature of the electrolyte.
For hole generating in steels, dilute acids are used but for the
form generating of steels aqueous solutions of inorganic salts such as
sodium choride, sodium nitrate or sodium chlorate are used either
singly or in combination.
Sodium chloride is widely used for ferrous materials due to its
superior machining rate and power efficiency. However, it produces an
inferior surface finish to that produced by a passivating electrolyte
82 such as sodium nitrate. In the case of a heat-resisting steel (X15
83 CrMolS), the metal removal/current density relationship was the same
for both sodium chloride and sodium nitrate but the surface finish
(Ra) after ECM at a current density of 0.60 A/mm2 was 2pm and 0.50pm
for sodium chloride and sodium nitrate respectively. A surface finish
of from 0.1 to 0.5pm Ra was obtained on other superalloy compositions
using a 20% sodium nitrate solution.
The surface finish obtained by the use of individual electrolytes
may be improved by the use of mixed electrolytes. For example,
certain alloys such as Nimonic 80A may be satisfactorily machined
using 10 to 20% sodium chloride solution. However, Nimonic 115 will
have an inferior surface finish but this may be improved by incor-
79 porating carbonate or phosphate ions. Similarly, the addition of 5%
sodium nitrate to a 25% sodium chloride solution also improved the
84 surface finish of Nimonic 115.
- 27 -
A high electrolyte conductivity enables BCM to be conducted at a
high current density using practical operating gap sizes and moderate
72 voltages. However, an improved surface finish may not necessarily
be obtained.
The specific conductivity of sodium chloride increases with
increasing concentration up to saturation (about 30%), ' 80 but an
iinproved surface finish is obtained with more dilute solutions. ItOC
has been shown, that for carbon steels and 18Cr8Ni stainless steel
an improved surface finish was obtained by decreasing the
concentration of sodium chloride from 20% to 5%. For a given feed
rate the gap distance required for ECM increases with increasing
concentration of electrolyte which accounts for the inferior surface
finish. 86
The electrolyte conductivity also increases with increasing
temperature such that sodium chloride solutions are 100% more
78 conductive at 71°C. than at 24°C. However, whilst the anode
efficiency is increased, the surface finish may deteriorate.
A reduction in the conductivity of an electrolyte of up to 50%
may occur as the result of the evolution of hydrogen gas. A
hydrogen bubble layer has been observed next to the cathode which
87 varies in thickness along the gap. Therefore, the local current and
hence the local rate of material removal will also vary thus causing a
* f - • u 88 variation in surface finish.
The "throwing power" of an electrolyte is an important factor in
controlling etching and pitting in inaccessible areas of the anode.
- 28 -
In contrast to electrodeposition a low throwing power is desired in
an ECM electrolyte in order to reduce the deleterious effect of
stray currents. The throwing power of sodium chlorate is much lower
than that of sodium chloride or sodium nitrate. When sodium
chlorate is used, the machining rate is faster than that of sodium
chloride because of its greater solubility and a surface finish of
81 0.05 to 0.125(jm was obtained on hardened steel. The low throwing
power of sodium chlorate is considered to be due to the passivation
of the anode in local current density regions. This was confirmed
when two passivating agents, bentriazol and potassium dichromate89 were added to sodium chloride. The throwing power of the sodium
chloride was reduced to the same value as sodium chlorate which
improved the dimensional control and surface finish on the
workpiece.
2.5.5 WORKPIECE MATERIAL
The surface finish produced by ECM is related to the type of
alloy, its microstructure and homogeneity.
Some superalloys are now being produced as single crystals but
metals are generally polycrystalline. Grain boundaries possess a
higher free energy than the crystal and are therefore preferentially
84 attacked during BCM to about 0.013 mm. A maximum grain boundary
penetration of 0.025 nm is generally considered an acceptable level90 for critically stressed parts. However, certain superalloys are
subject to excessive grain boundary attack in selective
electrolytes. For example, with certain Nimonic alloys, sodium
chloride tends to remove the grain boundaries preferentially but
- 29 -
with the addition of sodium nitrate, the grain boundary attack is
progressively reduced to a point where they are no longer attacked.
Correct adjustment of the chloride/nitrate mixture therefore results
in an acceptable surface finish. Similarly, carbonate and
phosphate ions when added to sodium chloride have the same effect. 79
The microstructure of pure metals and annealed solid solution
alloys consists of a single phase and a fairly uniform dissolution
results. Most alloys have duplex microstructures with phases having
different corrosion potentials so that preferential attack of some
phases will occur.
The microstructure of hypo-eutectoid carbon steels consists of
ferrite and pearlite. During corrosion in dilute acids the
91 cementite acts cathodically and the ferrite anodically, therefore
many localised corrosion cells are formed at the surface, which
increases the corrosion rate. Also, conducting metal sulphides have
a low hydrogen overvoltage so that manganese sulphide ( X = 0.10
ohm" cm" ) in steel act as local cathodes and initiate corrosion
92 attack. Evidence has been presented of the presence of active and
non-active sulphides in contact with a 3% sodium chloride solutions.
Active sulphides are surrounded by a fine dispersion of sulphide
particles which stimulates anodic dissolution of the matrix
92 resulting in the removal of inclusions and severe pitting.
With increasing carbon content more pearlite is formed which
also causes an increased localised corrosion rate in dilute acids.
This occurs due to the rapid corrosion of anodic areas when the
85 anode area is small compared with the cathode. Ito et.al. have
shown that this also applies to the BCM of steel in a sodium
- 30 -
chloride solution. For a given feed rate the surface finish
deteriorated with increase in carbon content from 0.19% to 0.52%.
This was attributed to the increased selective dissolution of
ferrite with increasing pearlite and decreasing ferrite content,
i.e. anode size effect.
For a steel with a given carbon content, the surface finish is
improved with decreasing grain size. The grain size of a normalised
steel is smaller than that of a similar steel in the annealed
condition, and the improvement in surface finish after normalising
is shown in table 2.5.
TABLE 2.5
REFERENCE
76
ii
n
82
it
n
MATERIAL STEEL
0.45%C.
n
n
C45
tl
11
HEAT TREATMENT
ANNEALED
NORMALISED
HARDENED
ANNEALED
NORMALISED
HARDENED
ELECTROLYTE
SODIUM CHLORIDE
M n
n n
SODIUM CHLORIDE
n n
n n
CURRENT DENSITY A/cm2
50
It
II
35
n
n
SURFACE FINISH
tun
3.5 Ra
2.0 "
1.0 "
29 Rt
11 "
2.5 "
A more comprehensive study has been conducted by Pramanik
93 et.al., using laboratory equipment where machining parameters
were carefully controlled. The electrolyte was a 6% w/v sodium
chloride solution, electrolyte flow rate 20.7 m/sec, electrode gap
- 31 -
0.4 mm and current density 71.7 A/cm2 . A variation in grain size
was determined by the ASTM method. The results also included the
effect of quenching and tempering on alloy steel. Table 2.6
TABLE 2.6
MATERIAL
ENSC. 0.35%
EN36C. 0.18%
Ni 3.52%
Cr 0.74%
HEAT TREATMENT
°C
NORMALISED 1010970940890840
ANNEALED 850NORMALISED 870HARDENED 860 TEMPERED 660HARDENED 860HARDENED 860 TEMPERED 550
ASTM GRAIN SIZE
X100
2-33-4
11 - 121616
X50034
-
-
5
METAL DISSOLUTION
RATE gm/min/annp
0.01540.01480.01370.01330.0127
0.01460.0139
0.0126
0.0106
0.0148
SURFACE FINISH
Uin
3.753.151.851.751.75
5.004.05
2.81
2.80
2.66
The results for the ENS steel show the improvement in surface
finish and decreasing dissolution rate with decreasing grain size.
94 Kbps and Quach. suggest that a small grain size should give an
increased dissolution rate, which is contrary to the above results.
The surface finish of steel EN36 after hardening is superior7fi op
to that after annealing and normalising as previously noted, ' '
table 2.6. However, it is significant that the surface finish
after tempering at 550 °C. is superior to that after tempering at
650°C. A difference in the corrosion rate in 1% H2S04 of a 0.95%
carbon steel after quenching and tempering has been shown by Heyn
and Bauer. 95 The corrosion rate increased with tempering
temperature up to 400 °C. and then decreased to the same value as
- 32 -
96 that of pearlite at about 700 °C. Uhlig has related this
behaviour to the changes which occur during the tempering of
martensite as described by Lament Averbach and Cohen. 97 During
tempering, low temperature martensite and epsilon carbide is
replaced by ferrite and cementite between 230°C. and 315°C. This
corresponds to the approximate tempering temperature of 300 to
95 400 °C. established by Heyn and Bauer for the maximum corrosion
rate. The superior surface finish produced by tempering at 650°C.
compared with tempering at 200 °C or to that of untempered
martensite has also been shown
electrochemical grinding (EGG).
98 martensite has also been shown by Geva et.al., both for ECM and
2.5.6 MECHANICAL PROPERTIES
Unlike EDM which drastically alters the surface structure,
properly conducted ECM is far less severe on the surface integrity
99 of the workpiece. It has been conclusively shown that ECM has
little effect on the YS, UTS, % elongation or %RA of steels and
super alloys. However, ECM removes surface compressive stresses
and leaves the surface in a stress free condition, which has the
effect of reducing the fatigue strength. In the case of a
Nimonic 80A machined in 15% sodium chloride, the removal of 0.20 mm
by ECM reduced the surface compressive stress from 600 MN/m2 to
zero. 102 The fatigue limit at 10 cycles was correspondingly
reduced from 340 MN/m2 to 260MN/m2 . Similar results have been
given by Rowden for low alloy and stainless steel and by
Gurklis103 for a H-ll alloy steel.
- 33 -
The use of an incompatible electrolyte and/or unsuitable
machining parameters may produce intergranular attack and/or "pitting"
in certain alloys. 104 ' 105 Evans et.al. 102 have shown that grain
boundary attack in a Nimonic 80A to a depth of 10~2 to 10~3 nm and 4 3 pitting to a depth of 10 to 10 mm had little further effect on the
fatigue life at 10 cycles. However, others have reported100 ' 103
that integranular attack to a depth of 0.013 mm in nickel based
superalloys reduced the fatigue life by 15%.
The fatigue life of alloys after ECM may be increased by lightly
cold working the surface. The fatigue life at 3 x 10 of a type 304
stainless steel was increased from 350 MN/m2 to 465 MN/m2 after vapour
blasting and to 510 MN/m2 after glass bead blasting. The comparable
fatigue strength after mechanical polishing to 0.4|jm surface finish
was 470 MN/m2 . Similarly the fatigue life at 10 cycles of a
Nimonic 80A was increased from 150MN/m2 after ECM to 355 MN/m2 after
shot blasting.
It has been stated that surfaces generated by ECM have better
wear, friction, corrosion and oxidation resistance than mechanically
finished surfaces. However, specific evidence is sparse. A
reduction in wear, coefficient of friction and running-in time of a
4.0%Ni2%Crl%W alloy steel has been reported. Spectrographic
analysis of the surface layers after ECM showed an increase in
chromium and nickel compared with mechanically machined surfaces. In
contrast, a reduction in surface hardness or "rebinder effect" occurs
in some alloys. An improvement in the corrosion resistance of some
materials after ECM has been reported but not extensively
investigated.
- 34 -
Cast alloys are generally less homogeneous than corresponding
wrought alloys and may therefore be expected to have an inferior
surface finish after ECM due to the selective dissolution of the
microstructure. For this reason a stainless steel casting would be
expected to have a surface finish of about 1.5pm compared with 0.1 to
90 1.0pm for its wrought counterpart. No evidence seems to be
available for the corresponding mechanical properties of cast metals
after ECM which forms part of the present investigation.
- 35-
CHAPTER III
3.0 EXPERIMENTAL PROCEDURE
3.1 MATERIALS
Both wrought and cast steels were used in the investigation. The
chemical composition was determined spectrographically by two
independent sources and the average composition is given in Table 3.1.
TABLE 3.1
IDENTITY TYPE
A
B
C
D
CAST
CAST
WROUGHT
WROUGHT
COMPOSITION (Wt %)
C
0.37
0.26
0.48
0.30
Si
0.32
0.40
0.26
0.34
S
0.023
0.018
0.044
0.013
P
0.023
0.027
0.032
0.014
Mn
0.66
0.94
0.85
0.70
Ni
0.26
1.32
0.12
2.80
Cr
0.27
1.62
0.10
0.56
Mo
0.09
0.41
<0.01
0.53
Cu
0.16
0.11
0.15
0.17
Sn
0.019
0.011
<0.01
0.018
The wrought steels were supplied as 20 mm. diameter hot-rolled
bars which were cut into suitable lengths. The cast steels were
produced commercially by the basic electric arc process using the
- 36 ~
conventional double-slag procedure. The final deoxidation was by
means of a furnace addition of ferro silicon, followed by 0.1%
aluminium added in stick form to the ladle during pouring from the
furnace.
The cast steel specimens were obtained from blocks 250 mm. long
and 100 mm square, cast in the conventional manner, Fig. 3.1. After
the removal of the risers, the blocks were sectioned by mechanical
sawing into 20 mm. square bars, either longitudinally, Fig. 3.2 or
transversely, Fig. 3.3
3.2 HEAT TREATMENT
The carbon steels A and C were normalised by heating to 920 °C for
one hour followed by cooling in still air.
The alloy steels B and D were heated to 880°C for one hour
followed by oil quenching. When sufficiently cool, the specimens were
immersed in liquid nitrogen in order to ensure the maximum
109 transformation of austenite to martensite. The specimens were
subsequently tempered at 600°C prior to further machining.
3.3 HOT ISOSTATIC PRESSING
Specimens were machined to 12 mm. diameter and hot isostatically
pressed using gaseous argon and ASEA Stora equipment . The HIP
parameters are given in table 3.2
- 37 -
TABLE 3.2
TEMPERATURE °C
930
1100
1160
1210
PRESSURE MN/m2
103
103
103
140
TIME HRS.
4
2
4
2
After HIP'ing, the specimens were re-heat treated and machined to
a smaller diameter before testing, to ensure the absence of any
surface connected porosity. The fatigue specimens were machined from
the initial 12.0 mm. to 3.81 mm. diameter prior to testing.
3.4 ELEiCTROCHEMICAL MACHINING
The final 0.25 mm. of some of the fatigue specimens was removed
by electrochemical machining using a Herbert-Anocut 150 machine. A
10% solution of sodium nitrate was used having a pH of 7.8 and an-1 -1
electrical conductivity (Re) of 0.10 ohm cm operating at a
temperature of 38 °C. The tooling consisted of a static cell which is
shown in fig. 3.4, located in the work area of the Herbert-Anocut
machine. The split copper electrodes and the finished fatigue
specimen are shown in fig 3.5. Specimens were machined at an inlet
pressure of 6.5 bar and at the maximum obtainable voltage of 15 volts
and a lower voltage of 10 volts. The machining parameters are given
in table 3.3
- 38 -
TABLE 3.3
STEEL
A + C
B + D
A + C
B
D
VOLTAGE volts
15
15
10
10
10
Current AmpsINITIAL
260
280
160
140
140
FINAL
180
180
100
100
100
Current DensityINITIAL A/cm3
77.2
83.2
47.5
41.6
41.6
FINAL A/cm3
53.5
53.8
29.7
29.7
29.7
MACHINING TIME
MDJS.
7-8
7-9
17
18 - 20
29 - 31
3.5.0 MECHANICAL PROPERTIES
3.5.1 TENSILE AND RELATED PROPERTIES
Tensile tests were carried out in accordance with BS18:Part2,
1971 using proportional round specimens having a gauge length of
5.65 So. The 0.2% PS,UTS, %EL and %RA were also determined.
3.5.2 NOTCH TOUGHNESS
Standard 10 mm square Charpy V-notch specimens were tested at
room temperature in accordance with BS131: Part 2, 1972. The average
value of a minimum of four specimens was taken.
- 39 -
3.5.3 FATIGUE TESTS
Pound specimens were carefully prepared by turning as
described in BS3518, Part 2, 1962. This was followed by the
introduction of successively finer scratches in a longitudinal
direction using emery papers of increasing fineness attached to a
Ludicke machine. The final polishing was carried out by the use of
felt "bobs" attached to a rotating spindle which were impregnated
with 6 pin and 1 pm diamond paste. This procedure ensured a final
surface finish of 0.06 to 0.10 pm Pa which is less than the maximum
of 0.127 urn stipulated in BS3518, Part 2, 1962.
The fatigue tests were carried out using rotating bending
machines with a zero mean stress (Sm = 0) rotating at 50 Hz. The
Wohler cantilever type (Avery 7304) was designed to use 6.68mm
diameter specimens whilst the Rolls-Royce machine used 3.81mm
diameter specimens. Some tests were conducted using notched
specimens with a theoretical stress concentration factor (K)
varying from 1.2 to 2.2. The form and dimensions of the specimens
are given in Fig. 3.6.
The results are expressed in the form of stress v log. N
graphs. Selected examples of the best fitting curves for the
finite life portion of the graphs obtained by computer are given in
Appendix I.
The S/N curves were derived using a package available at the
Polytechnic Computer Centre.
- 40 -
A weighted least squares polynomial is calculated by
Forsythe's method using orthogonal polynomials. The number of data
points (X,Y coordinates) must be at least 2 greater than the
maximum order polynomial.
The package proved to be too sophisticated for analysing the
data and a lack of time prevented further development of other
computation methods of curve fitting.
Two Rolls-Royce design machines were used, one was reserved
for the HIP experiments and the other for the ECM work. Similar
results were obtained when the same material was tested using both
machines, fig. 3.7.
For cast steel B after HIP at 1160°C, a statistical test was
conducted at two stress levels. The results of the Weibull
distribution and related statistical information are given in
Appendix II.
3.5.4 SHOT PEENING
After ECM, some of the fatigue specimens were shot peened in
order to introduce compressive stresses into the surface layers.
The shot peening was conducted using a stationary compressed air,
hand operated gun while the specimen was rotated in a small lathe.
The shot peening parameters are given in table 3.4
- 41 -
TABLE 3.4
PARAMETER
TYPE OF SHOT.
SHOT SIZE.
NOZZLE DIAMETER
DISTANCE FROM SPECIMEN
COMPRESSED AIR PRESSURE
SPEED OF SPECIMEN ROTATION
PEENING TIME
DETAILS
STEEL. HARDNESS 400 - 520 EL
BS2451, grade S120, mesh size
0.30 to 0.60 ran.
6 mm.
75 ran.
5.5. bars
0.58 Hz
4.0 minutes.
An Almen test, was carried out using "N" strips and the
saturation graph is given in fig. 3.8.
3.5.5 METALLOGRAPHY
Both the initial microstructures and those after subsequent
processing were examined by optical microscopy and where
appropriate by means of the scanning electron microscope (SEM).
The pearlitic steels A and C were etched in 3% nital and the
quenched and tempered steels B and D in Villela's reagent.
- 42 -
CHAPTER IV
Results
4.0.0 HOT ISOSTATIC PRESSING
4.1.0 TENSILE STRENGTH AND RELATED PROPERTIES
The tensile properties of cast steel A and B are shown in
Tables 4.1 to 4.5. Specimens taken from the edge of the casting are
differentiated from those taken from the mid and centre positions,
Figs. 3.2 and 3.3.
TABLE 4.1
MATERIAL - CAST STEEL A.
SPECIMENS - DIAMETER 7.98mm. GAUGE LENGTH 40 mm
POSITION
EDGE
MID/CENTRE
TREATMENT
NORMALISED 920°C.
NORMALISED 920°C.
0.2PSMN/m2
411 417
393 386
UTS MN/m2
673 673
679 674
EL.%
21.6 21.6
9.6 13.2
RA. %
22.8 23.2
14.7 15.4
TABLE 4.2
MATERIAL - CAST ALLOY STEEL B.
SPECIMENS - DIAMETER 5.64 mm. GAUGE LENGTH 28 mm
POSITION
EDGE
MID
CENTRE
TREATMENT
OQ 880°C. TEMP. 600°C.
OQ 880°C. TEMP. 600°C.
OQ 880°C. TEMP. 600°C.
0.2PS MN/m2
882 806
882857
878 863
UTS MN/m2
1018 1006
994 976
986 942
ELONG. %
15.5 11.2
4.3 5.0
4.4 3.2
RA. %
26.0 22.3
4.0 3.4
5.0 1.2
ENERGY TO FRACTURE J
88
11
10.5
- 43 -
TABLE 4.3
MATERIAL - CAST ALLOY STEEL B.
SPECIMENS DIAMETER 5.64 ran. GAUGE LENGTH 28mn
POSITION EDGE
TREATMENT
OQ.880°C. TEMP. 600°C.
HIP. 930°C. OQ.880°C. TEMP. 600°C.
HIP. 1100°C. OQ. 880°C. TEMP. 600°C
HIP. 1160°C. OQ. 880°C. TEMP. 600°C.
0.2PS MN/m2
882 806
999 999
913 906
999 995
UTS MN/m2
1018 1006
1100 1101
1026 1030
1101 1085
ELONG. %
15.5 11.2
15.86 14.15
15.7 16.1
5.53 6.75
RA. %
26.0 22.3
49.7 47.2
32.0 35.0
52.2 49.7
ENERGY TO FRACTURE J
88
110 93
143
40 48
TABLE 4.4
MATERIAL - CAST ALLOY STEEL B.
SPECIMENS DIAMETER 5.64itm. GAUGE LENGTH 28mm.
POSITION MID-POSITION
TREATMENT
OQ. 880°C. TEMP. 600°C.
HIP. 930 °C. OQ. 880°C.
TEMP. 600°C.
HIP. 1100°C. OQ. 880°C.
TEMP. 600°C.
0.2PS MN/m2
882 857
995 992
890 890
UTS MN/m2
994 976
1085 1078
1028 1025
ELONG. %
4.3 5.0
13.4 12.2
15.8 14.3
RA. %
4.0 3.4
49.7 52.2
28.0 26.0
ENERGY TO FRACTURE J
11
91 90
100
- 44 -
TABLE 4.5
MATERIAL - CAST ALLOY STEEL B.
SPECIMENS - DIAMETER 5.64 mm. GAUGE LENGTH 28 mm.
POSITION CENTRE
TREATMENT
OQ 880°C. TEMP. 600°C.
HIP. 1100°C. OQ. 880°C.
TEMP. 600°C.
0.2 PSMN/m2
878 863
910 864
UTS MN/m2
986 942
1026 1007
ELONG. %
4.4 3.2
9.5 14.5
RA. %
5.0 1.2
23.0 27.2
ENERGY TO FRACTURE J
10.5
50
4.2.0 NOTCH TOUGHNESS
The Charpy V-notch values of cast steel A and B are shown in
Tables 4.6 and 4.7 respectively.
TABLE 4.6
MATERIAL - CAST STEEL A.
POSITION
EDGE
MID-CENTRE
TREATMENT
NORMALISED 920 °C.
NORMALISED 920 °C.
CHARPY V-NOTCH J
32, 30, 32, 31
19, 20, 21, 22
AVERAGE J
31
20
- 45 -
TABLE 4.7
MATERIAL - CAST STEEL B.
POSITION
EDGE
KDGR
MID-CENTRE
MID-CENTRE
TREATMENT
OQ. 880°C. TEMP. 600°C.
HIP. 1100°C. OQ. 880°C.
TEMP. 600°C.
OQ. 880°C. TEMP. 600°C.
HIP. 1100°C. OQ. 880°C. TEMP. 600°C.
CHARPY V-NOTCH J
31, 32, 35, 40, 44, 46
48, 45, 47, 46
38, 25, 28, 29, 36, 24
40, 43, 40, 36
AVERAGE J
38
47
30
40
4.3.0 FATIGUE PROPERTIES
The fatigue limit of cast steel A before and after HIP is
given in Tables 4.8 and 4.9. The Kf value is calculated on the
basis of the fatigue limit of the edge specimens and also after
HIP. The corresponding S-N curves are given in Figs. 4.1 to 4.3.
TABLE 4.8
MATERIAL - CAST STEEL A. 6.68 DIAMETER SPECIMENS
POSITION
EDGE
MID-CENTRE
MID-CENTRE
EDGENOTCHED(Kt 2.2)
TREATMENT
NORMALISED920°C.
NORMALISED920°C.
HIP 1100 °C+NORMALISED
920°C.
NORMALISED920°C.
EL. MN/m2
240
215
300
154
FR.
0.36
0.32
0.45
0.23
Su,EDGE
—
1.12
—
1.56
Kf. HIP
1.25
1.40
-
1.94
q
-
-
-
0.62
FL % REDUCTION
-
10.4
-
35.8
FL % INCREASE
-
-
39.5
-
- 46 -
TABLE 4.9
MATERIAL - CAST STEEL A. 3.81mm DIAMETER SPECIMENS
POSITION
EDGE
EDGE
MID/CENTRE
MID/CENTRE
MID/CENTRE
TREATMENT
NORMALISED 920°C.
HIP 930°C.+ NORMALISED
920°C.
NORMALISED 920 °C.
HIP 930 °C. NORMALISED
920°C.
HIP 1100°C. NORMALISED
920 °C.
FL. MN/m2
300
345
245
300
310
FR.
0.45
0.52
0.37
0.45
0.46
K
-
~
1.22
~
^
H£P1.15
—
1.41
—
™
FL % REDUCTION
-
~
18.3
—
^
FL % INCREASE
-
15.0
-
22.4
26.5
The effect of prior homogenisation at 1100°C. on the fatigue
properties of edge and centre specimens of cast steel A is shown in
Table 4.10 and the S - N curves in Figs. 4.4 and 4.5.
TABLE 4.10
MATERIAL - CAST STEEL A. 3.81 ran. DIAMETER SPECIMENS
POSITION
EDGE
EDGE
MID/CENTRE
MID/CENTRE
TREATMENT
NORMALISED 920 °C.
ANNEALED 1100°C. NORMALISED 920 °C.
NORMALISED 920°C.
ANNEALED 1100°C. NORMALISED 920 °C.
FL. MN/m2
300
305
245
225
FR.
0.45
0.45
0.36
0.33
EDGE
-
-
1.22
1.33
FL %
REDUCTION
-
—
18.3
25.0
- 47 -
The fatigue properties of 6.68rom. diameter specimens of cast
steel B are shown in Table 4.11 and the S - N curves in Fig. 4.6.
For comparison, a series of specimens were prepared with machined
notches of known Kfc values. The results are shown in Table 4.12
and Fig. 4.7.
TABLE 4.11
MATERIAL - CAST STEEL B. 6.68 Itm DIAMETER SPECIMENS
POSITION
EDGE
MID/CENTRE
TREATMENT
OQ 880°C TEMP. 600°C.
OQ. 880°C. TEMP. 600°C.
FL MN/m2
385
220
FR.
0.385
0.220
K
-
1.75
FL %
REDUCTION
-
38
TABLE 4.12
MATERIAL - CAST STEEL B. 6.68 inn DIAMETER SPECIMENS
POSITION
EDGE
ii
ii
11
ii
TREATMENT
OQ. 880°C. TEMP. 600°C.
ii ii
ii ii
ii ii
ii ii
Kt1.0
1.2
1.4
1.8
2.2
FL MN/m2
385
355
325
250
220
FR
0.385
0.355
0.325
0.250
0.220
K
-
1.08
1.18
1.54
1.75
q
-
0.40
0.45
0.67
0.63
- 48 -
The variation in the fatigue properties with increasing
distance from the mould/metal interface is shown in Table 4.13 and
Fig. 4.8
TABLE 4.13
MATERIAL - CAST STEEL B. 3.81 mm DIAMETER SPECIMENS
POSITION
EDGE
MID
CENTRE
TREATMENT
OQ. 880°C. TEMP. 600°C.
ii ii
H ii
FL MN/m2
390
360
320
FR
0.39
0.36
0.32
EDGE
-
1.08
1.22
FL %
REDUCTION
-
7.7
18.0
The effect of HIP at increasing temperatures is shown in
Tables 4.14 to 4.16 and the S - N curves in Figs 4.9 to 4.14.
TABLE 4.14
MATERIAL - CAST STEEL B. 3.18 mm DIAMETER SPECIMENS
POSITION
EDGE
EDGE
EDGE
EDGE
TREATMENT
OQ. 880°C. TEMP. 600°C.
HIP 930 °C. OQ. 880°C. TEMP. 600°C.
HIP. 1100°C. OQ. 880°C. TEMP. 600 °C
HIP. 1160°C OQ. 880°C. TEMP. 600°C.
FL MN/m2
390
560
575
560
FR
0.39
0.56
0.575
0.56
Kf. HIP
1.44
-
-
-
FL %
INCREASE
-
44
47
44
- 49 -
TABLE 4.15
MATERIAL - CAST STEEL B. 3.81 mm DIAMETER SPECIMENS
POSITION
MID
MID
MID
MID
MID
TREATMENT
OQ. 880°C. TEMP. 600°C.
HIP 930 °C. OQ. 880°C. TEMP. 600°C.
HIP 1100°C. OQ. 880°C. TEMP. 600°C.
HIP. 1160°C. OQ. 880°C. TEMP. 600°C.
HIP 1210°C. OQ. 880°C. TEMP. 600°C.
FL MN/m2
360
530
540
560
550
FR
0.360
0.530
0.540
0.560
0.55
Hi!
1.51
-
-
-
-
FL %
INCREASE
-
47.0
50.0
53.0
55.5
TABLE 4.16
MATERIAL - CAST STEEL B. 3.81 ntn DIAMETER SPECIMENS
POSITION
CENTRE
CENTRE
TREATMENT
OQ. 880°C. TEMP. 600°C.
HIP 1100°C. OQ. 880°C TEMP. 600°C.
FL MN/m2
320
550
FR
0.32
0.575
H£P
1.80
-
FL %
INCREASE
-
71.9
To determine the possible contribution to the improvement
in mechanical properties which may be attributed to the
heating cycle used in the HIP treatment, specimens were
annealed at the HIP temperatures for the same time. The
results for cast steel B are shown in Tables 4.17 and 4.18 and
Figs. 4.15 to 4.18.
- 50 -
TABLE 4.17
MATERIAL - CAST STEEL B. 3.81 HTQ DIAMETER SPECIMENS
POSITION
EDGE
EDGE
EDGE
EDGE
TREATMENT
OQ. 880°C. TEMP. 600°C.
ANNEALED 1100°C. OQ.880°C. TEMP.600°C.
ANNEALED 1160°C. OQ.880°C. TEMP.600°C.
ANNEALED 1210°C. OQ.880°C. TEMP.600°C.
FL MN/m2
390
430
390
420
FR
0.39
0.43
0.39
0.42
FL %
INCREASE
-
10.2
NIL
7.7
TABLE 4.18
MATERIAL - CAST STEEL B. 3.81 nm DIAMETER SPECIMENS
POSITION
MID
MID
TREATMENT
OQ.880°C. TEMP.600°C.
ANNEALED 1160°C OQ.880°C. TEMP.600°C.
FL MN/m2
320
360
FR
0.32
0.36
FL %
INCREASE
-
12.5
4.4.0 MACROSCOPIC EXAMINATION
The macrostructure of cast steel A and B shows the presence of
columnar crystals at the periphery of the casting, Fig. 4.19. The
degree of porosity in the centre of the casting is not
representative since it is a section taken near to the riser.
- 51 -
4.5.0 MICROSCOPIC EXAMINATION
Unetched edge specimens of cast steel A and B showed type III
sulphide non-metallic inclusions with associated oxides, Fig. 4.20.
No microporosity was observed in edge specimens.
Specimens taken from the mid and centre of the casting
revealed the presence of randomly dispersed microporosity, Fig.
4.21. However, after HIP no evidence of microporosity was found in
these sections, but non-metallic inclusions showed signs of slight
spheroidisation Fig. 4.22.
After etching cast steel A in 3% nital, the as-cast structure
consisted of large crystals with a typical Widmanstatten ferrite
pattern, Fig. 4.23. After normalising, a fine grained ferrite
/pearlite structure was obtained, Fig. 4.24. The structure of the
specimens after HIP at 930°C. is shown in Fig. 4.25 and after HIP
at 1100°C. and 1160°C. in Fig. 4.26 which has a Widmanstatten
structure typical of an overheated steel. After subsequent
normalising at 920 °C. the structure of the HIP'ed specimens was
similar to that of Fig. 4.24.
After etching cast steel B in Vilella's reagent, the as-cast
microstructure showed evidence of microsegregation Fig. 4.27. The
structure after oil quenching from 880°C. and tempering at 600 °C
showed fine precipitated carbides Fig. 4.28.
The specimens from the mid and centre of the casting (steel B)
also showed microporosity which was not evident after HIP as in the
case of cast steel A. However a coarse microstructure was obtained
- 52 -
after HIP at successively higher temperatures, Figs. 4.29 to 4.34.
This was replaced by the normal microstructure after subsequent oil
quenching and tempering at 600°C., Fig. 4.28.
SEM photographs of typical fatigue fractures representing
edge, mid and centre specimens of cast steel B both before and
after HIP, are shown in Figs. 4.33 to 4.37.
- 53 -
CHAPTER V
RESULTS
5.0.0 MACHINING
5.1.0 MICROSCOPIC EXAMINATION
The basic microstructure of cast steel A and B is shown in
Figs. 4.24 and 4.28, and that of wrought steels C and D in Figs.
5.1 and 5.2. Selected optical photomicrographs of electro-
chemically machined surfaces taken by means of a Normarski
interference contrast objective are shown in Figs. 5.3 to 5.6.
Furthermore, representative SEM photographs of ECM surfaces of the
wrought and cast steels are shown in Figs. 5.7 to 5.14.
5.2.0 SURFACE FINISH MEASUREMENTS
Surface finish measurements (Ra) in the circumferential
direction are given in Table 5.1.
TABLE 5.1
SURFACE PREPARATION
MECHANICALLY POLISHED
ECM - 15 volts
ECM - 10 volts
WROUGHT STEELS Ra [iia.
0.002 0.003
1.00 1.05
2.00 3.00
CAST STEELS Ra pm
0.002 0.002
1.375 1.125
2.620 3.750
- 54 -
5.3.0 TENSILE STRENGTH AND RELATED PROPERTIES
The tensile strength and related properties of the base materials
are given in Table 5.2. Only edge specimens of the cast steels
were used.
TABLE 5.2
STEEL
A - CAST
B - CAST
C - WROUGHT
D - WROUGHT
HEAT TREATMENT
NORMALISED 920 °C.
OQ.880°C. TEMP.600°C.
NORMALISED 920 °C.
OQ.880°C. TEMP.600°C.
0.2 PS MN/m2
414
844
404
973
UTS MN/m2
673
1012
723
1085
EL %
21.6
13.4
23.2
15.4
RA %
23.0
24.0
48.4
58.0
5.4.0 NOTCH TOUGHNESS
The Charpy V-notch values of the base materials are given in
Table 5.3.
TABLE 5.3
STEEL
A - CAST
B - CAST
C - WROUGHT
D - WROUGHT
HEAT TREATMENT
NORMALISED 920 °C.
OQ.880°C. TEMP.600°C.
NORMALISED 920 °C.
OQ.880°C. TEMP.600°C.
CHARPY V-NOTCH J.
32, 30, 32, 31
31, 32, 35, 44, 46, 40
30, 30, 33, 34, 34, 36
86, 88, 90, 95, 97, 102
AVERAGE J.
31
38
33
93
- 55 -
5.5.0 FATIGUE PROPERTIES
The fatigue strength of the cast steels after mechanical
polishing, electrochemical machining and subsequent shot peening is
shown in Tables 5.4 and 5.5 and Figs. 5.15 to 5.18. For
conparison, graphs showing the effect of stress relief annealing
are included.
TABLE 5.4
MATERIAL - CAST STEEL A. EDGE SPECIMENS 3.81 ran DIAMETER
TREATMENT - NORMALISED 900 °C.
SPECIMEN PREPARATION
MECHANICALLY POLISHED
MECHANICALLY POLISHED + VAC ANNEALED 600°C.
ECM 15 VOLTS
ECM 10 VOLTS
ECM 10V + 15V + SHOT PEENING
FL MN/m2
300
270
280
260
380
FR
0.45
0.40
0.42
0.39
0.56
APPARENT Kf
-
1.10
1.07
1.15
-
FL %
DECREASE
-
10.0
6.7
13.3
-
FL %
INCREASE
-
-
-
-
46.2
- 56 -
TABLE 5.5
MATERIAL - CAST STEEL B. EDGE SPECIMENS 3.81 itm DIAMETER
TREATMENT - OQ.800°C. TEMP. 600°C.
SPECIMEN PREPARATION
MECHANICALLY POLISHED
MECHANICALLY POLISHED -1- VAC ANNEALED 600°C.
ECM 15 VOLTS
ECM 10 VOLTS
ECM 10V + 15V + SHOT PEENING
FL MN/ma
390
390
380
350
480
FR
0.39
0.39
0.38
0.35
0.48
APPARENT Kf
-
-
1.03
1.11
-
FL %
DECREASE
-
NIL
2.3
10.4
-
FL %
INCREASE
-
-
-
-
37.1
The ECM of cast steel B after HIP at 1100°C. is shown in Table
5.6 and Figs. 5.19 and 5.20.
TABLE 5.6
MATERIAL - CAST STEEL - EDGE SPECIMENS 3. Slum DIAMETER
TREATMENT - HIP 100°C. + OQ.880°C. + TEMP. 600°C.
SPECIMEN PREPARATION
MECHANICALLY POLISHED
MECHANICALLY POLISHED + VAC ANNEALED 600°C.
ECM 15 VOLTS
FL MN/m2
560
510
500
FR
0.56
0.51
0.50
APPARENT Kf
-
1.10
1.12
FL %
DECREASE
-
8.9
10.7
FL %
INCREASE
-
—
-
- 57 -
The effect of BCM and subsequent shot peening on the fatigue
strength of wrought steel C and D is shown in Tables 5.7 and 5.8
and Figs. 5.21 to 5.24.
TABLE 5.7
MATERIAL - WROUGHT STEEL C. SPECIMENS 3.81 ran DIAMETER.
TREATMENT - NORMALISED 920 °C.
SPECIMEN PREPARATION
MECHANICALLY POLISHED
MECHANICALLY POLISHED + VAC ANNEALED 600°C.
ECM 15 VOLTS
ECM 10 VOLTS
ECM 10V + 15V -1- SHOT PEENING
FL MN/m2
350
320
310
300
400
FR
0.48
0.44
0.43
0.41
0.55
APPARENT Kf
-
1.09
1.12
1.17
-
FL %
DECREASE
-
8.6
11.4
14.3
-
FL %
INCREASE
-
-
-
-
33.3
TABLE 5.8
MATERIAL - WROUGHT STEEL D. SPECIMENS 3.81 ran DIAMETER
TREATMENT - OQ.880°C. TEMP. 600°C.
SPECIMEN PREPARATION
MECHANICALLY POLISHED
MECHANICALLY POLISHED + VAC ANNEALED 600°C.
ECM 15 VOLTS
ECM 10 VOLTS
ECM 10V + 15V + SHOT PEENING
FL MN/m2
560
515
500
480
615
FR
0.52
0.47
0.46
0.44
0.57
APPARENT Kf
-
1.09
1.12
1.17
—
FL %
DECREASE
-
8.0
10.7
14.3
—
FL %
INCREASE
-
-
-
-
28.1
- 58 -
CHAPTER VI
DISCUSSION
6.0.0 HOT ISOSTATIC PRESSING
6.1.0 MICROPOROSITY IN CAST STEEL
The microscopic examination of the cast steel section (400mm x
400mm) used in the present investigation, revealed the existence of
random porosity at a distance greater than ~ 25mm from the
mould/metal interface. By its cuspidal morphology, it is
identified as interdendritic shrinkage porosity, Fig. 4.21. This
is caused mainly by the inability of the liquid metal to "feed"
through the interdendritic spaces during solidication to
accommodate the volume contraction accompanying the phase change.
Microporosity is related to the solidification process and
therefore to the resultant macrostructure. Thicker sectioned
castings solidify rapidly at the mould/metal interface forming
columnar crystals. As solidification proceeds, random nucleation
occurs ahead of the columnar crystals with the formation of
equi-axed crystals, Fig 4.19. No microscopically detectable
microporosity has been found in specimens taken from the edge of
the casting having columnar crystals, but it is invariably
associated with the equi-axed crystals, Fig. 4.21. This has been2, 112 previously reported by a number of investigators.
- 59 -
6.2.0 ESTIMATION OF THE EFFECT OF MICROPOROSITY ON FATIGUE
STRENGTH
Microscopical methods for determining the volume fraction of
microporosity were considered unsatisfactory since all porosity
must be assumed to be spherical in shape, which is not the case for
interdendritic shrinkage porosity. Furthermore, density
measurements do not accurately reflect the morphology or114 distribution of the porosity
In view of the cuspidal morphology of the interdendritic
porosity, the root radius and orientation of the cavities with
respect to the applied stress is more important than volume
fraction or density measurements. Since the fatigue strength is
particularly notch sensitive, it was considered that a quantitative
evaluation of the effect of microporosity would be obtained by
comparing the fatigue strength of specimens containing surface
microporosity, with edge specimens having machined notches of known
K values. It is shown in Table 4.11 and Fig. 4.6 that for 6.68 mm
diameter specimens, surface microporosity reduced the fatigue limit
of steel B from 385 MN/m2 to 220 MN/m2 which is equivalent to an
apparent Kf of 1.75. The results for similar notched specimens
with different K values are given in Table 4.12, Fig. 4.7. When
these results are compared by superimposing Figs. 4.6 and 4.7, it
is seen that the finite life fatigue strength of specimens
containing microporosity varied between a K of 1.5 and 2.2, and
the fatigue limit at 5.0 XI0 cycles coincided at a Kfc of 2.2 and a
Kf of 1.7. Since steel castings generally contain design features
- 60 -
with a Kt of less than 1.4, it is clear that the inicroporosity
found in the equi-axed crystal region of steel B represents a
severe notch effect.
6.3.0 EFFECT OF MICROPROSITY ON MECHANICAL PROPERTIES
A comparison of Tables 4.1 and 4.2 shows that inicroporosity
has little effect on the 0.2PS and UTS of cast steel A with a
pearlitic structure, but tends to reduce these properties in the
low alloy steel B which has a tempered martensite structure. This
may be attributed to the greater notch sensitivity of higher
strength steels. ' In contrast, the %EL and %RA are
drastically reduced in both steels to a level below the minimum
requirements of BS1300:1976 grades A5 and BT3 respectively. This
is also shown by a marked reduction in the area under the uniaxial
load-extension curve, which is a measure of the energy to fracture,
and may be regarded as an indication of the toughness of the
material. The %EA is particularly sensitive to the presence of114inicroporosity. Owing to their cuspidal morphology, the inter-
dendritic porosity becomes centres of stress concentration which
appear as "crow's feet" markings on the parallel portion of tensile
specimens. The micrcporosity therefore results in early crack
initiation and rapid crack propagation by void coalescence after
the maximum stress is reached, resulting in reduced "necking" of
the specimen.
From the limited results available, Table 4.6 and 4.7,
micrcporosity appears to lower the Charpy V-notch values (CVN). It
is clear that any porosity at the root of the notch, or in the path
- 61 -
of the propagating crack would result in lower energy values.118 Pattyn found the lowest CVN in the equi-axed crystal region, and
a narked anisotropy in specimens with columnar crystals. However,
119 120 other investigators ' have found little variation in CVN
values of specimens with columnar or equi-axed crystals from the
same casting. It would therefore be difficult to arrive at a
definite conclusion without careful experimentation and extensive
testing.
It is clear from Tables 4.8, 4.9, 4.11 and 4.13, Figs 4.1,
4.2, 4.6 and 4.8 that microporosity reduces the fatigue strength of
cast steel. The fatigue values of 6.68mm specimens are lower than
that of 3.81mm specimens indicating a marked fatigue "size121 effect". It is also evident that microporosity has a greater
effect in reducing fatigue in the 6.68mm specimens (38%) than in
3.81mm specimens(18%) , Tables 4.11 and 4.13. The reduction in
fatigue properties with increasing section size has been previously
reported in conventional castings , and in unidirectionally
122 solidified castings. This indicates that it is difficult to
achieve uniform fatigue properties throughout thicker sectioned
castings by means of improved foundry techniques. However, this
may now be achieved by means of hot isostatic pressing.
6.4.0 REMOVAL OF MICROPOROSITY BY HOT ISOSTATIC PRESSING
Cast steels containing microporosity subjected to HIP at
temperatures ranging from 930°C to 1210°C are free from any
microscopically observable porosity Fig. 4.22. Similar results
have been obtained in other metals which are reviewed in Chapter
II.
- 62 -
A complete explanation of the mechanism of pore closure in
castings by HIP is not yet available. However, it is generally
considered that the elimination of micropores occurs by the
diffusion of vacancies from pore surfaces to grain boundaries which123 act as "sinks". Coble and Flemings have shown this to be
dependent on the volume fraction and distribution of the porosity
and the grain size of the metal. An expression was developed
for the time required for the elimination of pores under isostatic124 pressure which has been used by Basaran et.al for the HIP of an
unidirectionally solidified low alloy steel. It was found that
after HIP at 200MN/m2 /1038°C and 1260°C/lhour, no microscopically
detectable porosity was observed, Fig 6.1. Furthermore, Stevens
and Flewitt have also reported the elimination of internal
porosity in a cast nickel base superalloy after HIP at
138MN/m2/1180°C. They concluded that the kinetics of the process
were consistent with existing theories of pore closure based on
vacancy diffusion. The application of isostatic pressure at a
suitable temperature induces radial and tangential forces around
the pores which may be approximated by means of thick wall cylinder18 equations. Final closure is obtained by diffusion bonding of the
collapsed surfaces. Evidence of plastic deformation and
recrystallisation in the area of previously existing porosity has
been observed in face centred cubic superalloys in the form of
twinned crystals.
6.5.0 EFFECT OF HIP ON MECHANICAL PROPERTIES
It is shown in Tables 4.4 and 4.5 that the HIP of cast steel B
containing microporosity at temperatures ranging from 930°C. to
- 63 -
1100°C. resulted only in a modest increase in the 0.2PS and UTS.
However, the %EL and %RA were substantially increased. This is
also shown by the increased energy to fracture represented by the
area under the uniaxial tension load extension curve. It has been1 *}f\
shown that in wrought steel, the %RA or strain to fracture is
solely a function of the volume fraction of second phase particles
such as non metallic inclusions. After HIP, only slight
spheroidisation of inclusions was observed, Fig. 4.22, therefore it
is clear that the increase in %RA was primarily due to the removal
of microporosity. In contrast, only a modest increase in the CVN
values occurred after HIP. This difference in behaviour between
%RA and CVN appears to be related to the high rate of strain used
in the Charpy test and the relatively shallow notch acuity
(r= 0.25mm) of the Charpy specimen, so that most of the energy is
used in crack initiation. The lack of consistency between %RA and127 CVN has been noticed in other cases, but it is considered that
in quenched and tempered steels the %RA values are a more128 discriminating criterion of toughness.
In terms of linear fracture mechanics, toughness is related to
the critical value of the plane strain stress intensity factor Kic
at the onset of unstable crack growth. Whilst a linear relation
ship has been reported between 1C and the CVN upper shelf values-L(_*
of high strength steels, 129 ' 1 Brown and Srawley found no
relationship between K and %RA. However, since Kic represents
plane strain conditions and %RA is a measure of the plastic
deformation which occurs prior to fracture, it is more probable
that a relationship exists between %PA and a mixed mode fracture
criteria such as the Crack Opening Displacement (COD). Supportive
- 64 -
evidence in favour of this suggestion is that the COD values at
crack initiation are independent of specimen geometry and so may be
considered as analogous to the formation of a neck in a tensile132 specimen. Furthermore, the fracture mode in both %PA and COD is
largely governed by type, volume fraction and spacing of second
phase particles. ' It would therefore appear that a marked
increase in COD values should be obtained after HIP, and merits
further investigation.
A major advantage of the HIP of cast steel containing
microporosity is the increase in the fatigue strength and the
fatigue limit as shown in Tables 4.8, 4.9, 4.15, 4.16 and Figs. 4.1
to 4.3 and 4.9 to 4.14. The increase is more marked in the higher
strength alloy steel B (47% to 72%) compared with the pearlitic
steel A (22% to 25%). It will be recalled that the reduction in
fatigue strength due to microcavities was greater in the alloy
steel B due to its greater notch sensitivity. Therefore it is
clear that their removal has resulted in a correspondingly higher
fatigue strength. Similarly the greatest reduction in the fatigue
strength (18%) was obtained in the centre specimens where
centreline shrinkage may have occurred, consequently the greatest
increase (72%) was obtained after HIP. It is therefore clear that
the severest interdendritic cavities in the centre of the test
casting were closed by means of HIP. The present work indicates
that at an argon pressure of 103MN/m2 , a temperature of
930°C/4hours or 1100°C/2hours is sufficient to close all internal
porosity in both the pearlitic steel A and the quenched and
tempered alloy steel B. In all cases the fatigue life and the
fatigue limit were elevated to the level of comparable wrought
steels, Fig. 6.2.
- 65 -
It will be seen from Figs. 4.34 and 4.35 that fractured
fatigue specimens from the mid/centre of the casting showed
evidence of fatigue crack nucleation from areas of surface
porosity, whilst similar specimens after HIP showed no evidence of
surface porosity, Fig. 4.36 and 4.37. The removal of stress
concentration due to surface microporosity would therefore result
in an increased fatigue strength.
The superior fatigue strength of cast steel after HIP may be
explained in terms of the mechanism of fatigue failure proposed by
Forsyth, which was divided into Stage I, initiation and
microscopic crack growth (mode II), and Stage II macroscopic crack
propagation (mode I).
In cases where crack like defects exist such as in cast steel
containing microporosity, Fig. 4.21 and 4.35, virtually the whole
fatigue life is occupied in Stage II crack propagation and may be
quantified by means of linear fracture mechanics. The relationship
proposed by Paris has been widely used, but is only strictly
valid for the intermediate zone of the growth rate curve. The rate
of crack growth may be expressed by the equation:
da—— =C(AK)m dn
where N is the number of cycles, C a material constant and m an
exponent which generally varies from 2.0 to 4.0 and exceptionally
up to 10. 137 The alternating stress intensity AK is the
difference between the maximum and minimum values for each cycle
( A K = Kmax - Kmin) .
- 66 -
For crack growth to occur, AK must exceed a threshold value AKc*
It has been shown that for a given R value, AK and AK arec
largely independent of composition, yield stress and138 139 raicrostructure. ' Fracture toughness, COD and AK values
O
for a wide variety of cast steels have been determined, 1 ' 140 and
may be used in design calculations.
The removal of microcavities by means of HIP produces a unique
situation where cast steel free from crack like defects is
produced. Therefore the traditional approach based on S/N curves
is more appropriate in this case since it is largely a test of a
materials resistance to crack initiation and microcrack growth
(Forsyth Stage I). This will require the initation of fatigue
cracks at heterogeneous nucleation sites at a free surface.
It has been shown that fatigue cracks in quenched and tempered
steels are initiated at persistent slip band intrusions and
extrusions which frequently emanated from grain boundaries and non141 metallic inclusions. The subsequent growth of microcracks was
found to be greater in the as quenched condition, and to decrease
with tempering temperature up to 650 °C. Since all specimens of
alloy steel B were finally quenched and tempered at 600°C. prior to
testing, it is clear that the increase in fatigue strength of the
specimens subjected to HIP was primarily due to the removal of
micro-porosity. The %RA also increased after HIP and may also142 influence the Stage I fatigue process. Duckworth has shown that
in high strength steels, the fatigue limit is related to the1 /I O
product of UTS X %RA. Dubinski et.al, also found a similar
relationship after microalloying with rare earth metals. The
-67 -
fatigue strength of wrought steels was little affected but that of
a similar cast steel was increased. A correlation was found
between the fatigue limit and the product of the tensile strength
and the reduction in area.
6.6.0 EFFECT OF HOMOGENISATION ON MECHANICAL PROPERTIES
During HIP, the steel is subjected to a high temperature for a
specified period of time, therefore it was considered important to
know the contribution made by the possible hoirogenisation of the
matrix in the improvement of the mechanical properties. The effect
on the fatigue strength was selected for investigation.
It is shown in Tables 4.10, 4.17 and 4.18 Figs. 4.4, 4.5 and
4.15 to 4.18 that annealing at the appropriate temperatures for two
hours resulted in no change in the fatigue strength of cast steel
A, and only a slight increase in the case of cast steel B.
Prolonged heating at a high temperature can cause a gradual change123 124 in the morphology of pores. However, Basaran et. al have
shown that after heating at 1315°C for 13 hours, the volume
fraction of microporosity was only slightly reduced, but after HIP
at 1200 MN/m2 /1250°C/lhour, the microporosity was completely
removed, Fig. 6.1. In an unidirectionally solidified high strength
cast steel free from microporosity, annealing at 1315°C for 50
hours was required to improve the fatigue ratio. In the present
work the increase in the fatigue strength of cast steel B due to
honogenisation amounted to a maximum of 12%, whilst the increase
due to HIP was from 47% to 72%. It is therefore clear that the
major part of the increase in fatigue strength after HIP was due to
the absence of microporosity.
- 68 -
6.7.0 EFFECT OF HIP ON THE MECHANICAL PROPERTIES OF CAST STEEL
FREE FROM MICROPOROSITY
It has previously been shown (Appendix VII) that the fatigue
strength of centre specimens after HIP was greater than that of
normal specimens from the edge of the casting. Therefore in order
to investigate a possible notch effect, edge specimens from the
columnar crystal zone were also subjected to HIP. It will be seen
from Table 4.3 that both the 0.2 PS and UTS of edge specimens of
cast steel B are improved. Whilst the EL% is little affected, the
RA% and the energy to fracture values are substantially increased.
However, as in the case of mid/centre specimens, the CVN values are
only moderately increased, Table 4.7.
The fatigue limit of edge specimens of cast steel A after HIP
is increased by 15%, Table 4.9, Fig. 4.2. However, the fatigue
limit of edge specimens of the higher strength cast steel B was
increased by 44 to 47%, Table 4.14 and Figs. 4.9 to 4.11.
It has been shown in Figs. 4.22 and 4.33 that edge specimens
having columnar crystals were free from microscopically detectable
microporosity. Therefore the improvement in mechanical properties
after HIP cannot be due to the elimination of porosity as in the
case of specimens from the mid/centre of the casting. However,
whilst the columnar crystals are free from microporosity and have
superior mechnical properties compared with equi-axed crystals, they
have been shown to be anisotropic. The RA% and the fatigue
strength are lower when the crystals are perpendicular to the
applied stress. 118 ' 12° In the present work the casting was
- 69 -
sectioned as shown in Fig. 3.2, so that the columnar crystals were
in this direction. This is shown in Tables 4.9 and 4.14 to be
equivalent to an apparent notch effect (Kf ) of =1.15 in cast steel
A and »1.44 in cast steel B.
The mechanical properties of cast steels are generally
determined from = 25mm section "keel" blocks, having a
predominantly columnar crystal structure, Fig 6.3 Although
representing the superior part of a casting, the mechanical
properties are inferior to those of comparable wrought steels in
the longitudinal direction. The notch sensitivity of cast steel
is also lower than that of wrought steel which makes the surface
finish of cast steel fatigue specimens unimportant. These two
factors are now shown to be due to the inherent notch effect of the
columnar crystals which is removed during the HIP cycle. The
fatigue strength of edge specimens after HIP is raised to the same
level as that of comparable wrought steels Fig. 6.2. The isostatic
pressure applied at an elevated temperature during HIP is
equivalent to hot working and evidence of recrystallisation
(twinned crystals) has been observed in the area of pre-existing32 cavities in FCC alloys. The cast plus HIP material will be
completely isotropic whilst after conventional hot working, the
steel remains anisotropic due to the elongation of non metallic
inclusions. It is therefore clear that hot isostatic pressing
results in the total improvement of a casting.
6.8 ECONOMIC FACTORS
It is recognised that HIP is an additional post casting
operation. Whilst the extra cost may be justified in the case of
- 70 -
titanium and superalloy castings, this may not be so for low cast
alloy steels. However, HIP may be necessary in specific cases
where high integrity and reliability is of importance.145 Furthermore, recent evidence has shown that for certain low
alloy investment castings, HIP was used to improve the EL% and Izod
values to meet specification requirements. In such cases the
additional cost may be justified.
- 71 -
CHAPTER VII
DISCUSSION
7.0.0 ELEiCTt^ClCHEiygCAL MACHINING
7.1.0 EFFECT OF ECM ON SURFACE INTEGRITY
The surface finish produced by the electro-chemical machining
of fatigue specimens is inferior to that obtained by conventional
mechanical polishing, Table 5.1. For a given electrolyte the
surface finish of electrochemically machined surfaces is strongly
influenced by the current density. When machining at 15 volts,
which is the safe maximum capacity of the machine, the current
density varied from 83 A/cm2 initially to 53.5 A/cm2 due to the
increasing gap size. The current density is on the lower limit of
that generally used for ECM which usually varies from 50 to
150A/cm2 , producing a surface finish (Ra) of from 0.25 to 1.0pm in
NaCl solutions. 146 Reported values for a 3molar sodium nitrate
solution vary from 0.65 to 1.0 ym which is of the same order as
that produced in the wrought steels used in the present work.
The surface finish of the cast steels machined under similar
conditions was inferior to that of the wrought steels, Table 5.1.
This is generally attributed to the greater heterogeneity of cast84 metals producing differential electrolytic attack.
The surface finish deteriorated markedly when the specimens
were machined at 10 volts with a current density of 47A/cm2 to 29.7
- 72 -
A/cm2 . General etching of the surface occurred as evidenced by a
black film which formed on the surface of the specimens. Also
selective attack of the microconstituents occurred particularly of
the cast steels Figs 5.6, 5.10 and 5.14. During the ECM of
pearlitic steels the ferrite is anodic and is selectively attacked,
whilst the cementite is cathodically protected. This is evidently
accelerated at a low current density particularly in cast steel A
Fig. 5.6 where a dendritic structure is revealed.
The surface finish produced by ECM is also dependent upon the
carbon content which in the present work is typical of medium148 carbon steels. It has been shown that in carbon steels machined
in NaCl solutions, the rate of dissolution and the surface finish
decreases with increasing carbon content from 0.13% to 1.5%. This
is clearly related to the microstructure of the steels. As the
carbon content increases up to 0.83% so the amount of cementite
increases at the expense of the ferrite which is the anodic
component of the galvanic cell. Furthermore, as the cementite
increases so an "anode size" effect is produced where the rate of
corrosion increases as the area of the anode decreases relative to
that of the cathode. Low current efficiency and an inferior
surface finish has also been encountered in the ECM of steels149 containing from 0.99 to 1.26%C in a 20% solution of NACl.
In ECM, the heat treatment given to the steel also influences
the surface finish produced. In the present work the carbon steels76 A and D were normalised, but it has been shown that a similar
steel (0.45%C) machined in NaCl gave an inferior surface in the
annealed condition than after normalising. The highest surface
- 73 -
finish was obtained when a martensitic structure was produced after
water quenching. Similar results have been reported by other82 93 investigators. ' The tenpering temperature of hardened steels
also affects the response to ECM. Increasing the tenpering
temperature of a low alloy steel (ASTM 4340) from 232°C to 658°C
increased the machining rate and iirproved the surface finish using78 98 a NaCl solution and a NaN03 solution. To ensure an uniform
microstructure, steels B and D were both tempered at 600 °C. after
oil quenching from 880°C.
The choice of electrolyte also influences the surface finish.
Whilst sodium nitrate has a higher resistivity and a lower
machining rate than NaCl the surface finish obtained is82 83 superior. ' This is related to the character of the surface
149 films formed due to passivation. In the present work the
surface finish of a 0.48% carbon steel in the normalised condition
machined in a 10% solution of NaNO., at a current density of 77 to
53.5 A/cm2 was of the order of l.Ourn Ra. However, it has been
reported that a steel of similar carbon content and heat
treatment, machined in NaCl at a current density of 50A/cm2 ,
resulted in a surface finish of 2.0pm Ra. Heat resisting alloys
machined using NaNO, had a superior surface finish to NaCl using83 the same machining conditions.
The surface finish is also related to the electrolyte
conductivity which increases with increase in operating
temperature. For this reason ECM is conducted at the highest
temperature consistent with smooth operating conditions, which in
the present work was 38°C. In general the electrolyte conductivity
and therefore the metal removal rate increases with increase in
- 74 -
electrolyte concentration, but the highest surface finish is notQC
always obtained. I to et.al have shown that for plain carbon
steels the surface finish is improved by a decrease in
concentration of NaCl from 20% to 10%.
7.2.0 EFFECT OF ECM ON MECHANICAL PROPERTIES
Mechanical properties such as 0.2PS, UTS and %RA are not very
surface sensitive and are therefore little affected by ECM. 100
However, the fatigue strength of metals is particularly sensitive
to changes in surface condition and may therefore be used as a
means of monitoring the surface integrity of machined conponents.
It is shown in Tables 5.4 to 5.8 and Figs. 5.16, 5.18, 5.20,
5.22 and 5.24 that ECM results in a decrease in the fatigue
strength of steel when compared with specimens prepared in the
conventional way by mechanical polishing. However, even carefully
prepared mechanically polished specimens contain surface
compressive stresses which increase the fatigue strength.
Therefore, the true fatigue strength of steel is obtained either by
electropolishing, or by vacuum annealing after mechanical polishing
in order to relieve the compressive stresses. Electropolishing
also removes non metallic inclusions from the surface thus leaving
cavities which may subsequently act as centres of stress
concentration, resulting in a lower fatigue strength. Therefore,
vacuum annealing at 600°C. was preferred in this work.
It is shown in Tables 5.7 and 5.8 and Figs. 5.21 and 5.22 that
a reduction of about 8% occurred in the fatigue strength of wrought
steels C and D after vacuum annealing at 600 °C. This is clearly
- 75 -
due to the removal of surface compressive stresses which were
generated during mechanical polishing. After ECM at 15 volts, (the
highest current density) the reduction in the fatigue strength was
about 11%. It is therefore clear that the reduction in the fatigue
strength which accompanies ECM is mainly due to the absence of
compressive stresses in the surface layers. This has also been102 shown by Evans et.al. , for a Nimonic 80A alloy.
When ECM was conducted at 10 volts, (low current density) the
reduction in the fatigue strength increased to about 14% due mainly
to the selective etching of microconstituents and associated
pitting, Fig. 5.12. In practice, ECM at a low current density
mainly occurs at areas which are not directly parallel with the
cathode surface. These results therefore show that when stray
current machining occurs, the additional reduction in the fatigue
strength is small. This is particularly applicable to ECM using a
10% Na NO, solution which has better dimensional control than NaCl
solutions. Evans et.al., also found that when stray current
conditions were simulated during the ECM of Nimonic 80A in NaCl
solutions, the resulting etching and intergranular attack (IGA) had
little further effect on the fatigue strength.
As previously reported (Appendix VII), the reduction in the
fatigue strength as the result of the ECM of the cast steels at 15
volts, Tables 5,4 and 5.5, Figs. 5.16 and 5.18 is less than that
which occurred in the wrought steels particularly the cast alloy
steel B. This is contrary to expectation in view of the greater
heterogeneity and inferior surface finish of the cast steels, Table
5.1. This clearly indicates that the cast steels are less affected
by the surface integrity than the wrought steels. This confirms
- 76 -
the evidence presented by Evans et.al 6 that cast steels are less
notch sensitive than comparable wrought steels. They found that
whilst the fatigue limit of highly polished cast specimens was
about 20% lower than that of comparable wrought steels cut
longitudinal to the rolling direction, the notched fatigue limit
(Kt 2.2) was the same for both steels. In addition, cast steel
fatigue specimens were found to be less sensitive to surface finish
than wrought steels. The fatigue limit of highly polished and
lathe turned cast steel specimens was of the same order, but the
fatigue limit of lathe turned wrought steel specimens was reduced
by about 28% when compared with the highly polished specimens. It
is known ' that the low notch sensitivity of cast steels is
due to the anisotropy of the columnar crystals. However, it is now
shown that this is removed when the specimens are subjected to hot
isostatic pressing (HIP), Table 5.6, Fig. 5.20.
7.3 EFFECT OF SHOT PEEKING ECM SURFACES
It will be seen from Tables 5.4, 5.5, 5.7 and 5.8, Figs. 5.16,
5.18, 5.20, 5.22 and 5.24 that shot peening increases the fatigue
limit of electrochemically machined wrought and cast steels. Shot
peening introduces compressive stresses into the surface layers but
has little effect in increasing the fatigue limit of highly1 ^*?
polished surfaces. However, increases in fatigue limit of up to153
100% have been reported in steel with as-cast surfaces. In the
present work the fatigue limit of electrochemically machined
wrought steels was increased by 28 to 33% and cast steels by 37 to
46%. The shot peening conditions and the Alinen rise are shown in
Table 3.4 and Fig. 3.8. The Almen rise figures were obtained
- 77 -
using "N" strips which is representative of light peening
conditions. The considerable increase in the fatigue strength is
therefore significant. Vapour blasting, glass bead and shot
peening has also been applied to electrochemically machined 403
stainless steel and Nimonic 80A, ' in order to increase the154 fatigue strength. It has been shown that the shot peening of
steel increases the fatigue crack initiation stage but has a
comparatively minor influence on crack propagation.
- 78 -
CHAPTER VIII
CONCLUSIONS
8.1 HOT ISOSTATIC PRESSING
i) Microporosity occurs in cast steel as the result of
the solidification process and is generally associated
with the areas which have equi-axed crystals.
ii) For a steel section 400mm x 400mm cast in a sand
mould, the inicroporosity when present at the surface had
an apparent K value of between 1.6 and 2.2.
iii) Microporosity in cast steel had little effect on the
0.2PS and UTS but decreased the EL%, RA% and fatigue
strength.
iv) Internal inicroporosity in carbon and low alloy cast
steels may be effectively closed by hot isostatic
pressing at an argon pressure of 103MN/m2 and a
temperature of 930°C for 4 hours, or 1100°C for 2
hours.
v) The HIP of cast steels containing internal micro-
porosity had little effect on the 0.2PS and UTS but the
EL% and RA% were increased.
vi) A major advantage of the HIP of cast steel is the
- 79 -
considerable increase in the fatigue strength, fatigue
limit, and the fatigue ratio. The values attained are
equivalent to that of comparable wrought steels.
vii) Homogenisation annealing at temperatures ranging from
1100°C to 1210°C resulted in only a slight improvement
in the fatigue strength of the cast steels.
viii) The improvement in the mechanical properties of cast
steel after hot isostatic pressing is mainly due to the
closure of microporosity.
ix) The fatigue strength of edge sections of cast steel is
also increased after HIP due to the removal of the notch
effect of orientated columnar crystals by isostatic hot
working.
8.2 ELEXITROCHEMICAL MACHINING
i) The surface finish produced after electrochemical
machining in a 10% sodium nitrate solution is inferior to
that of a ground surface, but that obtained in the case
of the wrought steels was superior to that of comparable
cast steels.
ii) Both the surface finish and the resultant fatigue
strength is strongly dependent upon the current density
used.
iii) When electrochemically machined at the highest current
- 80 -
density of 83/53 amps/on2 , the fatigue strength of both
wrought and cast steels was reduced.
iv) In spite of their greater heterogeneity, the reduction
in the fatigue strength was less in the case of the cast
steels than in the wrought steels. This supports the
finding that cast steels are less notch sensitive than
comparable wrought steels.
v) Mechanically polished specimens have a higher fatigue
strength due to the production of surface compressive
stresses. These may be removed by sub critical annealing
with a consequent reduction in the fatigue strength.
vi) When electrochemically machined at an adequate current
density, the fatigue strength is of the sane order as
that of stress relieved mechanically polished specimens.
Therefore, it is evident that a surface free from
microcracks and conpressive stresses is produced by ECM.
vii) ECM at a lower current density of 42/30 amps/cm2
resulted in a greater reduction in the fatigue strength
due to the selective etching of the microconstituents.
viii) The fatigue strength of electrochemically machined
steels may be considerably increased by the
reintroduction of surface conpressive stresses by means
of surface treatments such as shot peening.
ix) Shot peened ECM. surfaces have a fatigue strength in
- 81 -
excess of the original steel after mechanical polishing.
x) The increase in the fatigue strength after shot
peening is greater in the case of the cast steels than
that of the wrought steels.
xi) The fatigue strength of the steels electrochemically
machined at a low current density was increased after
shot peening to the same level as those machined at a
higher current density. This clearly shows the benefit
of shot peening areas which have been subjected to "stray
current" machining.
8.3 FURTHER VTORK
i) The effect of Hot Isostatic Pressing on the fracture
toughness (COD) of cast steel should be investigated.
Full advantage of the process may then be used in the
designing of steel castings.
ii) The HIP of cast die steels such as grades Hll or Hi 2
could be a major factor in improving the mechanical
properties and die life of these steels. The optimum HIP
parameters and the resultant mechanical properties should
be investigated.
iii) A detailed study of the optimum shot peening
conditions for the maximum improvement in the fatigue
strength of electrochemically machined steels would be
beneficial.
- 82 -
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APPENDIX I
COMPUTER PLOTS
550.
313
•fl CO
375
34
0
305
27
0
200.
PIG. Al.l
COMPUTER PLOT OP FATIGUE v
TIME.
REFER TO FI
G.
Cast
ste
el A
03
.81
mmEd
ge s
pecim
ens
- H.I.P
930
°C &
nor
mal
ised
920
°C•
Cent
re
M «
" "
' a
Edge
spe
cimen
s - n
orm
alis
ed 9
20°C
x
Cent
re
» «
«
TIM
E(H
OU
RS
) [
NU
MB
ER O
F R
EVER
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(N
) =
TIM
E (H
OU
RS
)»1.
8 x
10s
]
LOS
10
PIG. A1.2
COMPUTER PLOT OP FATIGUE v
TIME.
515:
>
Ji 41
a;
£ 37
5^U
J )
a: £
340J
305
Z7B
23
5
ID
Cast
stee
l B
06.6
8 mm
+ E
DGE
D M
ID
POSI
TIO
N
123
TIM
E (H
OUR
S)
U NU
MBE
R OF
REV
ERSA
LS (
N)
= TI
ME
(HO
UR
S)»1
.8 x
10s
]
. 4.
6
O.Q. &
Tem
p. 60
0 °C
LD£10
650
515.
560.
J
510|
,b %
4754
UJ o: oo
4403
UJ
370.
335
300
PIG. A1.3
COMPUTER PLOT OP FATIGUE v
TIME.
REFER
TO FIG. 4.8
Cast
ste
el B
O.
Q. &
Tem
p. 60
0 °C
<t>
3.81
mmED
GE P
OSIT
ION
MID
" CE
NTRE
H
' '
'4
...._
(HO
URS)
[
NUM
BER
OF R
EVER
SALS
(N
) =
TIM
E (H
OU
RS)
x 1.8
x10
5 3
650
580.
51 ^
4404
370
335
30
0
FIG.
A1.4
COMPUTER PLOT OF FATIGUE v
TIME.
REFER TO FI
G.
'l.K
J
Cast
ste
el B
03
.81m
mo
EDGE
SPE
CIM
ENS
OQ &
TEM
P 60
0°C
+ ED
GE S
PECI
MEN
S H.
I.P 1
100°
C +
OQ.-»
_
__
__
__
__
__
TE
MP
600
°C
123
TIME
(HOU
RS)
C NU
MBER
OF
REVE
RSAL
S (N
) = T
IME
(HOU
RS)x
1.8
x105
LOP1
0
t_n
650,
-£
51
B|
c/i
to UJ
CL
ID44
0
PIG. A1.5
COMPUTER PLOT OF FATIGUE v
TIME.
REFER TO
FIG.
4.
Cast
ste
el B
Ed
ge s
pecim
ens
03.8
1mm
+ H.
I.P 1
160 °
C Oj
Q.+
TEMP
600
°C
° O.
Q.+ T
EMP
600 °
C
TIME
C
NUMB
ER O
f RE
VERS
ALS
IN)»
TIM
E (H
OUR
S)-1
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W*
405:
37B
335
30
0
D^
0
1 •
' • '""1
' ' '""i
0 •
- '4
LOE1
0
S80
_
31
3:
-f. 1
4-10
3 +i B co
3
73
'L
U cc. 1
305
Z7O 233
200
FIG. A1.6
COMPUTER PLOT OF FATIGUE v
TIME.
REFER TO FIG. 5.16
Cast
ste
el A
Edg
e sp
ecim
ens
03.81
mm
Norm
alise
d92
0 °C
Edge
spe
cimen
s E.
C.M.
15
Volt
+ ii
O
il
II II"
10
"»
10&1
5"
+ sh
otpe
eniny
0
TIM
E(HO
URS)
C
NUM
BER
OF R
EVER
SALS
(N
) =
TIM
E (H
OU
RS)
x 1.8
x 1
0s
LOG
T0
61
3;
sea,
54
5
51
0:
UJ ot: H- tyi
405
37B
333
PI'7.
A1.7
COMPUTER PL
OT OP FATIGUE v
TIME.
REFER TO
FI
G. 5.
18
CAST
STE
EL B
- ED
GE S
PECI
MEN
S O.
Q 60
0°C
* TE
MP
600°
C4-
ECM
10&
1S V
OLT
' SHO
T PE
ENIN
G
• E
CM 1
5 VO
LT
O EC
M 10
VOL
T
TIME
(HOU
RS)
[ NU
MBER
OF
REVE
RSAL
S (N
) = T
IME
(HOU
RS}x
1.8 x
10s
LOG
10
700.
663
585
CD
-z.
"£. •M g to
l/> UJ
sza.
PIG. A1.8
COMPUTER PLOT OP FATIGUE v
TIME.
REFER TO FIG. 5.20
CAST
STE
EL B
- ED
GG S
PECI
MEN
S O.
Q 88
0°C
« TE
MP
600°
C
• HI
P 11
00 °C
a
HIP
1100
°C»£
CM 1
5V+
PO
LISH
ED
3B3J
39B
J
TIM
E(HO
URS)
C
NUM
BER
OF R
EVER
SALS
(N
) = T
IME
(HO
URS)
x 1.8
x 1
0sLO
S TO
550
313 j
445;
to to £
30S
27
0
233
FIG
. A
1.9
C
OM
PUTE
R
PLO
T O
P F
AT
IGU
E
v T
IME
. R
EFER
TO
F
IG.
5.2
2
WRO
UGHT
STE
EL C
; NO
gMAL
ISgD
920
°£
__
__
__
__
_
"+"
ECM
10 &
15 V
OLT
* SH
OT P
EENJ
NG
•
ECM
15 V
OLT
O EC
M 10
VOL
T
TIME
(HOU
RS)
[ NU
MBER
OF
REVE
RSAL
S <N
) = T
IME
I HO
URS)
* 1.8
x10*
3
LOG
10
790.
713.
Ul
QC
373.
540.
47H
:
4OO
.
PIG
. A
1.10
CO
MPU
TER
PLO
T O
P FA
TIG
UE
v TI
ME.
RE
FER
TO
FIG
.
WRO
UGHT
STE
EL D
- O
.Q.8B
O°C
* TE
MP
600°
C
X EC
M 10
&15
VOIT
* S
HOT
PEEN
ING
O PO
LISH
ED
• E
CM 1
5 VO
LT
+ E
CM 1
0 VO
LT
TIM
E (H
OUR
S)
C NU
MBE
R OF
REV
ERSA
LS (
N)
= TI
ME
(HO
URS)
x 1
.8 x
10s ]
LOG
10
APPENDIX II
WEIBULL DISTRIBUTION
TABLE Al.l
FATIGUE TESTS AT A CONSTANT STRESS
FATIGUE STRESS MN/m2
CYCLES TO FAILURE (N)
SPECIMEN SEQUENCE
123456789
590 MN/m2
N
1.44 x 10c1.44 x 10J?1.62 x 10e1.62 x 10"1.80 x 10c1.80 x 10^2.00 x 10;?2.34 x lol?4.32 x 10
620 MN/m2
N
7.2 x 10*7.2 x 10"?9.0 x 10:9.0 x 10;9.0 x 10:1.1 x 10;1.1 x lot1.3 x 10^1.44 x 10s
TABLE A1.2
STATISTICAL RESULTS
SAMPLE SIZESHAPE <£>MEAN LIFE (N)STANDARD DEVIATION (a)
FATIGUE STRESS 590 MN/m2
93.7
2.04 x 10.8.48 x 10
FATIGUE STRESS 620 MN/m2
93.7 ,
1.00 x 10"2.33 x 10
The shape factor obtained from the Weibull analyses indicates a near normal distribution, for which the mean fatigue life (N) and the standard deviation is shown above, Table Al.2.
A2.1
0 = 620 MN/rn?
r::: S'::iDQ\TSPK!MENS:llJ£11QQlC1,1,00' '• .: FATIGUE-
2_^^^» « « ]" Vaan 49.5
3.7
AGE AT FAILURE
FIG. A2.1
WEIBULL DISTRIBUTION AT A STRESS OF 620MN/M2
A2.2
CU
MU
LATI
VE
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R C
ENT.
FA
ILU
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o o
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tr tn
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mr->
0
>
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m_<
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LH
vO
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O o
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APPENDIX III
EXPERIMENTAL PROCEDURE
FIG. 3.1
METHOD OF CASTING
A 3.1
25
C = Centre specimens -12=Peripheral specimens
Dimensions in mm
FIG. 3.2
SECTIONING OF CASTING - LONGITUDINALLY
XX)
E
M C M E
/ /
/
//
/
/s)
C= Cmtrt speciatns M= Mid. • E= Edg« •
Jfifl.
FIG. 3.3
SECTIONING OF CASTING - TRANSVEESELY
A 3.2
FIG 3.4a HERBERT - ANOCOT MACHINE
FIG. 3.4b STATIC rrrr.T. POSITIONED FOR MACHINING
A 3.3
FIG. 3.5
SPLIT COPPER ELECTRODES
A3.4
DIMENSIONS IN MILLIMETRFS
r<*>7 -0.012 r03.810 ±0.012 R38.1-
--r (a)
-012 nominal
R70
11.86
(b)
k-2.5
r$6.68
-—————-V- (0
^Specimen drawn showing notch form with smallest and largest radii
PIG. 3.6 FATIGUE SPECIMEN DIMENSIONS
A3.5
U)
380-
360-
340-
~ 3
20H
.0 CO
CO
300-
CO
280-
260-
240
Cas
t ste
el B
No
rmal
ised
920°
C
\•
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spe
cim
ens-
mac
hine
no.1
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——
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|—•
I 'I
i |—
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&
10s
• 10
6 10
7CY
CLES
TO
FAILU
RE (
N)
FIG
. 3.
7 FA
TIGU
E RE
SULT
S FR
OM D
IFFE
RENT
MAC
HINE
S
10B
> » -J
50-
40-
30-
ALM
EN"N
"(m
m)20
-
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2060
80
TI
ME
(SEC
S.)
100
120
FIG.
3. 8
GRAP
H SHOWING AL
MEN RI
SE
APPENDIX IV
RESULTS; HOT ISOSTATIC PRESSING
HU
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FIG
. 4.
1 CA
ST S
TEEL
A.
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CT O
F PO
ROSI
TY A
ND H
IP
AT 1
100°
C.
6.68
ran
DIA
. SP
ECIM
ENS
NJ
450-
400-
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300
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L
Cast
ste
el A
03
.81
mm
104
+ Ed
ge s
pecim
ens
-H.I.
P 93
0 °C
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orm
alis
ed 9
20°C
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ntre
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cimen
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orm
alis
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20 °C
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ntre
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12)
(2)
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-i • •
• • i 10
6 CY
CLES
TO
FAILU
RE
"L 10
710
8
FIG
. 4.
2CAST S
TEEL
A.
EFFECT OF PO
ROSI
TY AND
HIP
AT 9
30°C
. 3.81 i
tm DIA
. SP
ECIM
ENS
U)
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350-
300-
S
250-
200
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Cast
ste
el A
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Spec
imen
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m•
Cent
re s
pecim
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mal
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ntre
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Nor
mal
ised
o
Edge
spe
cimen
s -N
orm
alis
ed
10s
106
CYCL
ES T
O FA
ILUR
E (N
)
107
108
PIG
. 4.
3CAST STEEL A.
EFFECT OF POROSITY AND HIP
AT1100°C.
3.81 nm DIA.
SPECIMENS
380
360-
340-
D trt £
320H
g +
. 300
-
280-
260-
240
Cast
ste
el A
Ed
ge s
pecim
ens
0 3.8
1 mm
• Ann
eale
d 11
00 °C
+ No
rmali
sed
920°
C + N
orm
alise
d920
°C
• i 10s
106
CYCL
ES T
O FA
ILUR
E (N
)
• i 107
108
FIG.
4.4
CAST STE
EL A.
EFFECT OF ANNE
ALING AT
11
00 °C. 3.81 r
an DIA.
EDGE SPE
CIME
NS
U1
320-
300-
^
280-^
LU
fM
CC
^
CO ^
= «
260
-1ID
240-
220-
Cent
re
104
+ An
neal
ed 1
100
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orm
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20 C
•
Nor
mal
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°C
105
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CYCL
ES T
O FA
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„10
7
FIG
. 4.
5 CA
ST S
TEEL
A.
EFEE
CT O
F AN
NEAL
ING
AT
1100
°C.
3.81
irm
DIA
. CE
NTRE
SPE
CIM
ENS
->
10U
500-
Cas
t st
eelB
06
.68m
mO.
Q. &
Tem
p. 6
00 °C
400-
•H to
t/1 UJ a:
300-
200-
T '
'10
"10
5
FIG
. 4.
6
106
108
CYCL
ES TO FAILURE
(N)
CAST
STE
EL B
. EF
FECT
OF PO
ROSITY
6.68 mm DIA
. SPECIMENS.
500
Cas
t st
eel
B 0
6.66
mm
400-
+ 1 6 en U
J cc30
0-
O.Q.
& T
emp.
600
°C
200
-i- ED
GE
PLAI
N S
PECI
MEN
S D
X-
O E
DGE
NOTC
HED
SPEC
IMEN
S 3xPl
ain
Kt1.
2
Kf2.2
105
106
CYCL
ES T
O FA
ILUR
E (N
)10
fl
FIG. 4.7
CAST STEEL B.
EFFECT OF MACHINED NOTCHES
WITH DIFFERENT Kt VALUES.
6.68 i
tm DIA.
SPECIMENS
650-
600-
Cast
ste
el B
O.
Q. &
Tem
p. 60
0 °C
03
.81m
ma
EDGE
POS
ITIO
No
MID
«o
CENT
RE
••
550-
oo
<S 500-
UJ a: o
450-
400-
350-
300 10
"10
510
6 CY
CLES
TO
FAIL
URE
(N)
107
FIG. 4.8
CAST S
TEEL B
. EFFECT OF POROSITY 3
.81
ntn DIA. SPECIMENS
D ——
O ——
-o(2J-
-0(2)-
10
700-
>
650-
600-
b 55
0-
00
UJ ex.
500-
450-
400
350
i.i-
Cas
t st
eel B
Ed
ge p
ositi
on
03.8
1 mm
10*1
• H.
I.P 9
30 °C
.+ T
EMP
600
C 0
O.Q.
+ TE
MP
600
°C
10s
106
CYCL
ES T
O FA
ILUR
E (N
)10
7
___
FIG. 4.
9 CAST STEEL B
.EFFECT OF HIP AT 93
0°C.
3.81 mm DIA EDG
E SP
ECIM
ENS
1 (2) (2)
a • a-
650-
600-
550-
Cas
t st
eel
B 03
.81m
m°
EDGE
SPE
CIM
ENS
QQ. &
TEM
P 60
0°C
• ED
GE S
PECI
MEN
S H.
I.P 1
100°
C +
O.Q.
+ TE
MP
600°
C
K2J
-
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QC
CO
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450-
400-
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DO
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D
CD
D
..
350-
300-
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—,
, ,
, r,
——
——
——
,——
——
r10
6 CY
CLES
TO
FAIL
URE
(N)
10*
105
107
FIG
. 4.
10
CAST
STE
EL B
. E
FF
EC
T OF
HIP
AT
110
0°C
. 3.
81 D
IA.
EDGE
SPE
CIM
ENS
108
700-
Cas
t st
eel
B Ed
ge s
pecim
ens
03.8
1mm
650-
H.I.P
116
0 °C
OjQ.
+ TE
MP
600
°C
O.Q.
+ TE
MP
600
°C
^
2:
600-
b 55
0-to
oo U
J
500-
450-
400-
Bn
x-0
a
a a
a o • *
•(2)
350- 10
"-.—
——
.——
| .
i •
r |
——
——
——
.——
——
r-10
6 CY
CLES
TO
FAIL
URE
(N)
108
FIG. 4.11
CAST STEEL B
. EFFECT OF HIP
AT 1
160°C. 3.81 i
tm DIA.EDGE SPECIMENS
FATIGUE STRESS (a) ±MN/m2
O-iin
LHo
Oo
«J1uiunLHo
Oo
Os V/1O
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3 %•= m
JLCO
VD H
(—1t—m on
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i
>
z:
650
600-
550-
500-
at UJ ID
450-
400-
350
300
Cas
t ste
el B
O.
Q. &
Tem
p. 60
0 °C
03
.81m
m
\
t
10"
i -X
' •
•
\ o o
o o
° MI
D PO
SITI
ON•
" "
+H.I.
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o CE
NTRE
POS
ITIO
N»
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HIP
110
0 °C
105
106
CYCL
ES T
O FA
ILUR
E (N
)
107
FIG. 4.13
CAST STEEL B
. EFFECT OF HIP AT 1
100°C.
3.81 nm DIA. MID AND CENTRE SPECIMENS
••(2
)
10
650
600-
550-
b
500-
i/>
in UJ o:
400-
350
300
"\
N+
,."
\o
o
H.I.P
121
0°C
* O.
Q.+
TEM
P 60
0°C
H.I.P
116
0°0
QQ.+
TEM
P 60
0°C
O.Q.
+ TE
MP
600°
CC
ast
stee
l B
03.8
1 mm
Mid
-pos
ition
10"
10s
106
107
CYCL
ES T
O FA
ILUR
E (N
)FI
G.
4.14
CA
ST S
TEEL
B.
EFFE
CT O
F H
IP
AT 1
160°
C a
nd 1
210°
C 3
.81
nrm D
IA.
MID
PO
SITI
ON
SPE
CIM
ENS
108
650-
600-
550
2: co CO
LiJ
UJ
500-
450
400
350
300
Cas
t st
eel
B Ed
ge s
pecim
ens
03.8
1 mm
Anne
aled
110
0 °C
O.Q.
& Te
mp. 6
00 °C
'O
.Q. &
Tem
p. 60
0 °C
•\-
V
X
•
"(2)a a
—
I •
- I
• I
. .«
105
106
107
CYCLES T
O FAILURE
(N)
FIG. 4.15
CAST STEEL B
. EFFECT OF ANNEALING AT
1100°C.
3.81 rm DIA.
EDGE S
PECIMENS
108
650
600-
<£
550-
z: § 5
00H
LU
CL •
400H
350-
300
Cast
ste
el B
Ed
ge s
pecim
ens
03.8
1 mm
•X °•a
105
O.Q.
& T
emp.
600
°CAn
neal
ed 1
160
°C &
O.Q
.& T
emp.
600 °
C
-,—,—
J-, T_
r_,
——
—,—
T_
106
CYCL
ES T
O FA
ILUR
E (N
)
-rr—
T10
710
8
FIG
. 4.
16
CAST
STE
EL B
. EF
FECT
OF
ANNE
ALIN
G AT
11
60°C
. 3.
81 i
tm
EDGE
SPE
CIM
ENS.
650
600-
550-
| ;
§ 50
°"CO cc
. I—
l/l45
0-
400-
350-
300
Cast
ste
el B
Ed
ge s
peci
men
s 03
.81
mm•
Anne
aled
121
0 °C
& O
.Q.&
Tem
p. 60
0 °C
a O.
Q.&
Temp
. 600
°C
(3)
(2)
106
CYCL
ES T
O FA
ILUR
E (N
)
107
108
FIG. 4.17
CAST S
TEEL,'B.
EFFECT OF ANNEALING AT
1210°C.
3.18 i
rni DIA.EDGE S
PECIMENS
650
Cas
t st
eel
6 M
id po
sitio
n 03
.81
mm60
0-O.
Q.&
Temp
. 600
°CAn
neal
ed 1
160°
C &
O.Q.
& T
emp.
600
°C
550-
oo
£ 5
00-
UJ
QC45
0-Ll
J "- 4
00-
350-
a(2
)0(2
)'
300 10
4-i
——
i—r-
r-r-
105
106
CYCL
ES T
O FA
ILUR
E (N
J
10'
108
FIG
. 4.
18
CAST
{jT
KKb
B.
EFFE
CT O
F AN
NEAL
ING
AT
1160
°C.
3.81
nm
DIA
. M
ID P
OSI
TIO
N S
PECI
MEN
S
FIG. 4.19a
MACR3STRUCTURE CAST STEEL A
FIG. 4.19b
MACROSTRUCTURE CAST STEEL B
A4.19
»•• ». •
'V
FIG. 4.20a
NON METALLIC INCLUSIONS. X250
FIG. 4.20b
NON METALLIC INCLUSIONS. X1000
A4.20
FIG. 4.21a
MICROPOROSITY. MID POSITION X250
FIG. 4.21b
MICROPOROSITY. CENTRE XlOO
A4.21
FIG. 4.22a
CENTRE SPECIMEN AFTER HIP AT 1100°C. X200
FIG. 4.22b
CENTRE SPECIMEN AFTER HIP AT 1100°C. XlOOO
A4.22
OOIX 'Do 026*V T3SLS I.SVD
OOIX J,S\D SV 'V THSLS J.SVD
tt'W
OOIX 'DoOOII 3H dIH
oolx "Dooee iv din
FIG. 4.27
CAST STEEL B. AS CAST. X100
FIG. 4.28
CAST STEEL B. OIL QUENCHED 880°C AND TEMPERED 600°C. X500
A4.25
FIG. 4.29
CAST STEEL B. AFTER HIP AT 930 °C
FIG. 4.30
CAST STEEL B. AFTER HIP AT 1100°C.
A4.26
LZ'W
IW d •g T331S JLS\fD
Do09TI O^f dIH •g T33LS
FIG. 4.33a
CAST STEEL B. EDGE SPECIMEN. FATIGUE FRACTURE X20
FIG. 4.33b.
CAST STEEL B. EDGE SPECIMEN. FATIGUE FRACTURE X200
A4.28
FIG. 4.34a
CAST STEEL B. MID SPECIMEN. FATIGUE FRACTURE X20
FIG. 4.34b
CAST STEEL B. MID SPECIMEN. FATIGUE FRACTURE X200.
A4.29
FIG. 4.35a
CAST STEEL B. CENTRE SPECIMEN. FATIGUE FRACTURE X20
FIG. 4.35b
CAST STEEL B. CENTRE SPECIMEN. FATIGUE FRACTURE X200
A4.30
FIG. 4.36a CAST STEEL B.
CENTRE SPECIMEN AFTER HIP AT 930 °C. FATIGUE FRACTURE X200
FIG. 4.36b
CAST STEEL B.CENTRE SPECIMEN AFTER HIP AT 930°C.
FATIGUE FRACTURE X200
A4.31
FIG. 4.37a CAST STEEL B.
CENTRE SPECIMEN AFTER HIP AT 1100°C. FATIGUE FRACTURE X20
FIG. 4.37b.
CAST STEEL B.CENTRE SPECIMEN AFTER HIP AT 1100°C.
FATIGUE FRACTURE X200
A4.32
APPENDIX V
RESULTS: ELECTROCHEMICAL MACHINING
FIG. 5.1. WROUGHT STEEL C.
NORMALISED 920°e. x 200.
'f BBS.
FIG. 5.2. WROUGHT STEEL D.
OIL QUENCHED 880°C TEMPERED 600°C x 1000
A5.1
FIG. 5.3WROUGHT STEEL C.
BCM 15 VOLTS. X 500
FIG. 5.4. WROUGHT STEEL C.
BCM 10 VOLTS, x 500
A5.2
FIG. 5.5. CAST STEEL A.
BCM 15 VOLTS, x 500
FIG. 5.6. CAST STEEL A.
BCM 10 VOLTS. X 500
A5.3
FIG. 5.7a. WROUGHT STEEL C.
ECM 15 VOLTS, x 500
FIG. 5.7b. WROUGHT STKKI • C.
ECM 15 VOLTS, x 2000
A5.4
FIG. 5.8a.WROUGHT STEEL C.
ECM 10 VOLTS, x 500
FIG. 5.8b.WROUGHT STEEL C.
ECM 10 VOLTS. X 2000
A5.5
FIG. 5.9a. CAST STEEL A.
ECM 15 VOLTS, x 500
FIG. 5.9b. CAST STEEL A.
ECM 15 VOLTS, x 1000
A5.6
FIG. 5.10a. CAST STEEL A.
BCM 10 VOLTS, x 500
FIG. 5.10b. CAST STEEL A.
ECM 10 VOLTS, x 1500
A5.7
FIG. S.lla. WROUGHT STEEL D.
ECM 15 VOLTS, x 500
FIG. 5.lib. WROUGHT STEEL D.
ECM 15 VOLTS. X 2000
A5.8
FIG. 5.12a. WROUGHT STEEL D.
BCM 10 VOLTS.x 500
FIG. 5.12b. WROUGHT STEEL D.
ECM 10 VOLTS, x 2000
A5.9
FIG. 5.13a. CAST STEEL B.
BCM 15 VOLTS, x 200
FIG. 5.13b.CAST STEEL B.
ECM 15 VOLTS, x 1000
A5.10
FIG. 5.14a.CAST ffl'KKT. B.
ECM 10 VOLTS, x 1000
FIG. 5.14b. CAST STEEL B.
ECM 10 VOLTS, x 2000
A5.ll
to
450-
400
CO S35
0H
- CO L
U 10 5
300H
250-
200
_i—
——
i——
I—i—
* T-
*
104
03.8
1 mm
CA
ST S
TEEL
A
- NO
RMAL
ISED
920
°C
+ ED
GE S
PECI
MENS
- PO
LISH
EDo
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- POL
ISHE
D &
ANNE
ALED
600
°C
i i
i i
i 10s
1 i
r i
i [ 10
6—
a?
(2)
10CY
CLES
TO
FAILU
RE (
N)FI
G.
5.15
. CA
ST S
TEEL
A.
EFFE
CT O
F ST
RESS
REL
IEF
ANNE
ALIN
G.
500-
I .
. ,-
U,
J...
., ,1
J_Ca
st s
teel
A E
dge
spec
imen
s 03B
1 mm
Norm
alise
d °
\JJ
920
°C
N,
450H
Q -
NV°
O
400-
CO
CO
UJ oc CO
U
J ID
350-
300-
250-
200
o Ed
ge s
pecim
ens
E.C.
M. 1
5 Vo
lt +
„ ,.
M 10
,.
Q •>
« «
10&1
5"
+sho
tpe
ening
a QN
Q ax
a a
Q\a
aa
*
° 4-
>-f
(2)
a—
>
o——
^D
——
>
(2) .o
——
H
(2)if
—
10*
105
106
10CY
CLES
TO
FAIL
URE
(N)
FIG
. 5.
16.
CAST
STE
EL A
. EF
FECT
OF
ECM
AND
SHO
T PE
ENIN
G.
600
550-
soo-
•n to I/) E
a: j—
i/) UJ
450-
p 4
00-
350-
300
Cast
ste
el B
Ed
ge s
pecim
ens
03.81
mm
O.
Q.88
0°C+
Tem
p. 60
0 °C
Mech
anica
lly p
olish
edst
ress
rel
ief
600 °
C +—
+—
10"
105
106
107
108
CYCL
ES T
O FA
ILURE
(N)
FIG. 5.17.
CAST STEEL B
. EFFECT OF
STRESS R
ELIEF ANNEALING.
Ln
600
550-
500-
450-
cc
i— CO g400-
350
300
03.8
1mm
i i.ii.i_
__
•
. i
. .. .1
__
_
. .
i . .
i. i—
-S
D CA
ST S
TEEL
B -
EDGE
SPE
CIM
ENS
O.Q.
880°
C *
TEM
P 60
0°C
aao
ECM
10&1
5 VO
LT +
SHO
T PE
ENIN
G o
ECM
15 V
OLT
* E
CM 1
0 VO
LT
I '
' "l
c 10
51
•• I
' ' '' I
.10
6To
^
(2)
•a-
~o-
(2)
(2)
-i—|
i •
• i 10'
CYCL
ES T
O FA
ILUR
E IN
)
FIG
. 5.
18.
CAST
STE
EL B
. EF
FECT
OF
ECM
AND
SHOT
PEE
NING
.
650-
03.8
1 mm
600-
55°"
i— C/l
UJ
500-
450-
400-
350
CAST
STE
EL B
- ED
GE S
PECI
MENS
O.Q
.880
°C+
TEMP
600
°Cx
HIP
1100
°C -M
ECHA
NICA
LLY
POLI
SHED
o
HIP
1100
°C-
H «
+STR
ESS
RELI
EF60
0 °C
x x •o|2
)-
1 I
' '
"I,-
——
—'—
—'—
' I
' '
' 'I,—
——
'——
'—'
I '
' "I
,——
—'—
—'—
r~l
' '
'^5
106
107
1010
FIG
. 5.
19.
10s
CYCL
ES T
O FA
ILUR
E (N
)
CAST
STE
EL B
. H
IP A
T 11
00°C
. EF
FECT
OF
STRE
SS R
ELIE
F AN
NEAL
ING.
650-
_L03
.81m
m
600-
550-
z: •H LU cm
500-
3 4
50-
400-
CAST
STE
EL B
-EDG
E SP
ECIM
ENS
Ofl.
880°
C *
TEM
P 60
0°C
x HI
P 11
00°C
a H
IP11
00°C
*EC
M 1
5V•
POLI
SHED
• •
x X (2)
•o-
350-
T-—
—•—
• '
i ' '
"I.—
—•—
—'
' i
' ' "
i——
•—
105
106
107
CYCL
ES T
O FA
ILUR
E (N
) CA
ST S
TEEL
B.
HIP
AT
1100
°C.
EFFE
CT O
F EC
M.
104
FIG
. 5.
20.
10
151
500
450-
1/1 £350
-1
o:
300
250-
200
• i
• . .
i
0 3.8
1mm
WRO
UGHT
STE
EL C
- NO
RMAL
ISED
920
°C-P
OLI
SHED
-PO
LISH
ED &
ANN
EALE
D 60
0°C
J2)
io5
io6
io7
io8CY
CLES
TO
FAIL
URE
(N)
FIG. 5.21.
WROUGHT ST
EEL
C.
EFFECT OF
STRE
SS RELIEF ANNEALING.
500-
450-
400-
to CO f=
300-
250-
03.81
mm
WRO
UGHT
STE
EL C
- NO
RMAL
ISED
920
°CQ
- EC
M 10
&15V
OLT
* SH
OT P
EENI
NG
+ -
POL
ISHE
D o
- EC
M 15
VOL
T * -
ECM
10
VOLT
10b
10'
108
CYCL
ES T
O FA
ILUR
E (N
)
FIG
. 5.
22
WRO
UGHT
STE
EL C
. EF
FECT
OF
ECM
AND
SHO
T PE
ENIN
G.
to
O
700
"SI
+1 LU a:
600-
£5
00
400
—Li
iW
roug
ht s
teel
D
Spec
imen
s 03
.81 m
m O.
Q. 8
80 °C
+Te
mp.
600
°C
V \
o M
echa
nicall
y po
lishe
d-f-
«
« +
stre
ss r
elie
f 60
0 *C
~"~r
<'•"
!—
105
106
CYCL
ES T
O FA
ILURE
(N)
(2) 12).
107
"T"
108
FIG
. 5.
23.
WRO
UGHT
STE
EL D
. EF
FECT
OF
STRE
SS R
ELIE
F fiN
NEAL
ING.
NJ
750-
700-
O-J e f 6
50-
co £
600-
o: H
- to LLJ |
550-
LU
500-
*f 3
U
, ,
1 .
. .
, I
, .
. |
, ,
t,l
. ,
, 1
, .
. ,1
....
, ,,!,,,
* 3.
81m
m
WRO
UGHT
STE
EL D
- O
.Q.8
80°C
+ TE
MP
600°
C
Q EC
M 10
& 15
VOL
T +
SHOT
PEE
NING
\Q
Q
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LISH
ED
X
o £C
M 15
VOL
T X
*
ECM
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OLT
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S
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n "^
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a
O
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o-
/ '
'c
' ' i
' 'T
'
-
1010
10°
10'
CYCL
ES T
O FA
ILURE
(N)
10
FIG.
5.24.
WROUGHT
STEE
L D.
EF
FECT
OF EC
M AN
D SHOT PEENING.
APPENDIX VI
DISCUSSION
§1-H
CC
O o
0.5-
o
• AS
CAS
To
HT 1
3 HO
URS
AT 1
315
°C*
ii 30
n
" 13
15 °C
o
HIP
1 HO
UR A
T 12
50°C
& 2
9000
psi
(200
M P
a)
0.5
1DI
STAN
CE F
ROM
CHILL
(cm)
FIG. 6.1
MTCRDPORDSITY
IN CAST STEEL
AFTER HE
P (R
EF.
124)
•H
700-
600-
500-
400-
300-
200-
100-
- Cast Steels after HIP 1100 °C o Wrought Steels• Cast SteelsUnnotched-Wrought
Endurance Ratio 0-5
0-400
Unnotched-Cast 8
.xx.xxxx-
Notched-Cast Wrought
T600 800
U.T.S.1000 1200
FIG. 6.2 FATIGUE STRENGTH OF CAST AND WROUGHT STKH'.-S (REF. 116)
A6.2
FIG. 6.3 M^RQSTBDCTUKE OF CAST STEEL KEEL BLOCK (BEF. 114)
A6.3
APPENDIX VII
PUBLICATIONS
TECHNOLOGYA PUBLICATION OF THL ML I'ALS SOUL
SEPTEMBER 19812 VOLUME
Page 349 Effect of hot isostatic pressing on fatigueI. STRODE. ML ». BASSET, AMD C. DA VIES
355 Superplastic behaviour during compression of as-cast alloyO. A. MASSEF. M. SUEMY. AMD A. EL-ASHRAM j,
360 Influence of reheating temperature, 680*-1280*C, on extrusion of Aid']steel and low-C. high-Mn steel ->,*->*»»'»,K. E- HUQHEB. M. L. PLAVT. AMD C. M. SELLAMS
368 Influence of hot-working parameters on earing behaviour sheetT. SHEPPAMD AMD M. A. ZAIDI
375 Filler metals for containers holding irradiated fuel,P. M. MATHCW, F. WIIHSEM. J. S. MADCAM. AMO A. C. By
, 'V " V -'^
381 Technical note: Effect of titanium mocutotkut'^ impact strength of chromium-manganese aMoA. SASAK. J. PEHMIMO. AMD J. MLEWUM*
BISITS list page 385; tear-off contents facing full list of contents inside front cover
Effect of hot isostatic pressing on fatigue strength of cast steel
I. StrodeM B BassettC Davies
The inadvertent occurrence of porosity m cast \teel has tended in restrict the u.u 1 of lasting a\ a means o) producing component requiring a high degree of structural tntegrit\ Evidence is now presented showing that hoi isostatic pressing M an effective \\av of removing internal porosity, resulting in a marked improvement in mechanical properties. particularl\ the Jatigue strength. The deleter ton* effcri of tnterdendritic porositv on the mechanical properties of tii\i Meel i\ \hn\\n. together with the improvement in the fatigue strength after hot iMKitatic pressing at I UK) C at a pressure oj !4HMPa The independent tilect of temperature was determined h\- subjecting 'sound' cast steel specimens to a homogemzatton heat treatment at the same temperature and time interval that were used during hot isostatic pressing. The significant improvement in fatigue strength reported tna\ he due to a change in non-metallic inclusion morphology. The added tost oj the process mav he justified where a high degree of structural integr,t\ M imperative. MT-814
< /WO The Metals Society. Manuscript received 2 July IV8I Mr Strode and Or Bassett are in the Department of Mechanical and Production Engineering, and Dr Davies is in the Department of Chemical Engineering. The Polytechnic ofWales. Ponnpridd. Mid Glamorgan. Wales.
For mam >ears. much research has been directed towards producing high-integrity steel castings b> foundrv methods designed lo produce controlled directional solidification The success of these methods is evidenced hv the increasing use ol cast steel components for import am engineering applications '
External and internal defects in steel castings for highly -.tressed parts can usually be detected onl> by a high degree of non-destructive testing While external defects ma\ sometimes be rectified by judicious welding, the presence of internal defects often results m the rejection of the component with a consequent increase in production costs. However, hot isostatic pressing iHlPl now offers a new approach tu this problem.
In HIP the components are simultaneously heated and subjected to isostalic pressure by means of a liquid or gas. hence consolidating the cast material bv causing internal pores to close OngmalK. the process was developed for the diffusion bonding of atomic fuel elements, and later it was extended to the manufacture of sintered hard metal parts " Now it is beme used increasing!) to eliminate porosity in high-alloyed cast steels for gas-turbine and aerospace applications '
It has been shown"1 that the properties of comparable cast and wrought steels differ considerably, the cast steels having lower tensile and fatigue strengths This has been attributed to the as-cast dendritic structure, which in wrought steels is consolidated b\ mechanical working Owing to the nature of the solidification process, specimens from the edge of the casting having columnar crystals were free from micro- cavmes. while specimens taken from the central areas having equiaxed crystals contained interdendntic porosity Therefore, these two types of specimen were isolated and treated separately
The purpose of the investigation was to study the fatigue properties of specimens taken from the peripheral and central regions of a casting, in the normalized condition.
after homogemzation. and after HIP to close internal porosity
Although preliminary HIP experiments were conducted at a temperature of 1190 C. it was feared that this might have caused gram-boundary liquation Therefore, it was decided to use a lower and more economical temperature of IIOOC
Experimental details
MATERIAL AND HEAT TREATMENT The steel used was produced commercially by the basic electric arc process, using the conventional double slag procedure, and deoxidized using ferrosilicon and aluminium. The composition is given in Table 1. Test specimens were obtained from blocks 250mm long and 100mm square, cast in the conventional way (Fig.l). After the risers were removed, the blocks were sectioned longitudinally into four equal slices 22mm square, thus \ieldmg twelve edge specimens and four from the central area (Fig.2). These were given a normalizing treatment by heating to 920 C for I h, cooling in still air. and tempering at 600 C
MECHANICAL PROPERTIESTensile tests were earned out in accordance with BSI8: 1971. and Charpy V-notch tests were performed at room tem perature using 10 mm square specimens, as specified m BS 131 Pi2 1972.
Fatigue specimens were carefully prepared by turning as prescribed in BS35I8 . Pt 2 1962 The final polishing was earned out in the longitudinal direction using first 6 um and then 1 \im diamond paste. The faugue tests were initially carried out on a Wohler-type rotaung cantilever machine (Avery type 7304), using a 668mm dia. tapered specimen.
Table 1 Composition of steal, wt-%
S<
032
P
0 023
Metals Technology September 1982 Vol 9 349
350 Strode et at Effect of press.ng on fatigue strength
.-07-0012 -03810*0012
R 38.1-1
50.80
__ I38Z)
-0668
1 Method of casting test blocks
- ————— ---* (c)
After fracture, the ends of these specimens H 2-0 mm du.l were machined to 3 81 mm dia. and prepared for testing on a Rolls-Royce rotating cantilever machine. The specimen dimensions are shown in Fig. 3.
HOT ISOSTATIC PRESSINGThe 120mm dia broken fatigue specimens from the centre of the cast block were subjected to HIP using gaseous argon
* Specimen drawn showing notch form with smallest and largest radii
a Rolls Rovce type testpiece. b c Wohler-iype testpiece 3 Dimensions of machined specimens, in mm
112
11
10
2
C
c9
3
C
C
8
4
5
6
7 C = centre specimens 1-12 = edge specimens
and ASEA Slom equipment. The specimens were processed ji a temperature of 1100 C; the operating conditions are shown in Table 2. These specimens were subsequently normalized at 920 C. machined to 381 mm dia. (Fig.3). re- prepared, and tested on the Rolls-Royce rotating bending- faiigue machine
Results
MACROSTRUCTUREAn etched section taken from the upper end of the cast block is shown in Fig.4. The columnar crystals at the edge of the block are clearly visible, as are the equiaxed crystals in the central areas. The amount of microporosity evident is not
Table 2 Operating conditions for hot isostatic pressing
: Initial Heating i
«• —— — - ---- ~-P- Tempe«a(u'e, 'C 1100
2 Sectioning of test block to produce edge and centre T(me m|n 9Q
lo Process Cooling to
1 100 9001484 1320 120 25
Metals Technology September 1982 Vol 9
Strode ei at Effeci ot pressing on tongue strength 351
4 Mac restructure of ai-c*»t block
representative of the specimens. since the sccnon *as taken from the region tmmediaieh under the feeder head
MICROSTRUCTUREPolished specimens taken from the edge ot ihe casting were relatively sound, but dn example of the microporosilv found in specimens taken from the centre of the block is show n in Fig 5 After HIP treatment. specimens from the centr.il areas of the casting were found lo he free from microporosiiv ipn> 6l After normali/ing Ji ^2" (J j re)jti\el> nne-grjmed fernte pcjrlne Mruuurc is ohumed iFig^i Substantial gram growth occurred during HIP treatment at 1 100 C. resulting m a t>pical Widmansiaiten structure (FigKj This necessitated further norrruh/ini: !«> restore (he structure ii 1 ihat of the orijiiiul specimen-,
MECHANICAL PROPERTIESThe effect of microcavmes on the mechanical properties of centre specimens is shown in Table * The fjiituc limit Jl 5-(J » 10 c>cles of edge and centre specimens jfier normalising at 920 C. and after HIP treatment and subsequent normalizing, is gi^en in Tables 4 jnd > and
6 Non metallic inclusions m centre specimen after HIP treatment
sh.mn in KILS ID and I I Both cdj»f and centre specimens were gi\en a homoireni/ation annealing at I 100 C for 2 h to simulate the heating ocle jri\en during HIP treatment The suhscqufnt results arc gi^en in Table r> and shown in hits i: jnd I •
Discussion of results
hriim Table ^ a mas be seen that mterdendriiic p<irosit> had little effect <tn the proof stress and I TS. but reduced Ihe
5 Microporosity in untreated centre specimen7 Fernte-pearhte structure after normalizing at 920"C.
specimen etched in 3% mtal
Metals Technology September 1982 Vol 9
352 Strode ei at Effect of pressing on fatigue strength
• edge specimenso edge specimens
(notdxd,Ki=22)• centre specimens o centre specimens
8 Widmanstatten structure after HIP treatment at 1 lOO'C. specimen etched m 3% nital
10° 107 CYCLES TO FAILURE
10 Effect of HIP treatment on fatigue «tr 6 68mm dia specimens
ioP
ength.
elongation and RA values, as reported in previous work s " Although the Charpv V-notch values were also reduced, other investigators have reported conflicting results A comparison of Tables 4 and 5 and Figs 10 and II shows lhat the latigue limit decreases with increasing specimen diameter owing to the 'sue effect' This has been attributed to the different stress gradient or to the increasing probability of cavities occurring in large sections While a size effect is also apparent in wrought >teel specimens," it is less significant in cast steel specimens consolidated b> HIP treatment (Tables 4 and 5)
\ comparison of the mechanical properties of specimens taken from the edge and centre of a cast steel block shows clearly the reduction m the fatigue strength of the centre specimens, which is due to the presence of microcavmes Since the I'TS is little affected b> microcavines, the reduction m the fatigue ratio is particularly marked When the fatigue limit of the edge specimens is taken as the base
value, an indication of the notch effect of the microcavuies ma> he obtained b> companng the K, value of the centre specimens with that of artificially notched specimens with K, = 2 2 (Table 4 Fig 10) According to Frost ei al.* the K, and K, values approach each other at low hardness values. so that in this case K, = 1 14 = K, However, if the fatigue limn after HIP is taken as the base value, then K t = 1 24 for the centre specimens and 1-09 for the edge specimens.
During HIP treatment considerable austenite gram grow in occurs I Fig.8). which necessitates a further normalizing treatment to refine the grain structure (Fig.9). The low-c>cle and high-cycle fatigue strengths are both increased (Tables 3 and 4. Figs 10 and I 11 A fatigue limn of 300 MN m : is achieved (Table 4|. which compares favourably with that of longitudinal wrought steel specimen^ of similar diameter, composition, and heat
3 Structure after HIP treatment at 1 100'C and normal izing at 920'C: specimen etched in 3% nital
b~30O
^2bO-
o<20O
edge specimens. normalized • centre specimens, ncrmaiizad
centre specimens, normalized *HIP
re" TCYCLES TO FAILURE
•K?
11 Effect of HIP treatment on fatigue strength. 3 81 mm dia. specimens
Table 3 Effect of microcavitie* on mechanical properties
Heat treatment
Normalized at 920'C No<malaed at 920'C
Elongation, %Impact(Ch«rpv V-notch), J
41' 393
673679
21696
3232:30 192220
Metals Technology September 1982 Vol 9
Strode er al Effect ot pressing on fatigue strength 353
Table 4 Effect of hot isostatic pressing and normal tiing on fatigue limit of 6 68 mm dia. speci-
Pos.iian
Edye
Centre
Cenire
Edge- notched'«, 2 2)
•an 0
Faligue limit. F.mgue tdlio T'eaimenr MNm ; ( FL UTS) K t
Noinviiwed ai 240 0 36920'CNormdliiea ai 215 0 32 1)2
HIP and 300 0 45noimahzed at920 CNormalised ai 154 0 23 1 56920 C
\Z
*' Cb$ 260l~
1/1 260-
O5 2401-
2 2O 1 in5
o normalized at 92O°C
normaiizad at 92O°C
\o o\ *• \ o -|
0 O ^"V^
oCD
O O+~
Table 5 Effect of hot isostatic pressing and normal izmg on fatigue limit of 3 81 mm dia speci mens, stress mode. rotating bending SmMn 0
Position
Edge
Ceni.e
Centre
Fjugue iimiTreatment MNm •
Normalised al J85920 CNormaii/ed at 250920 CHIP and ' 310
"20 C
I Fatigue rat.oFL UTSi «,
042
037 114
046
treatment "* "' ll has been shuwn' ' in.u HIP irvaimenl h.ii little effect on the I TS. M> ihe impro\emcni in tjimuc ratio ii even arejier
An important feature of centre speumcn> alter HIP treatment is that the fjiiguc limn i-> greater ihan is ohumed m the normalh treated edge specimens While irm would suggest I he presence of minute cavities m the edge specimens, no such cutties were found during examination b\ microscope The improvement m the fatigue properties after HIP ireaiment mav therefore he due lo a combined conv>lidation and thermal efTecl
The HI P treat men l enables the horn ogeni/a( ion of microconstituents to occur, which is wh\ some specimens were subjected to a heat-treatment c>cle. without the application of pressure, designed to simulate the efTecis of HIP
Figure 12 shous that onl\ a shghl impro\ement in fatigue strength resulted from homogeni/mn specimens containing
CYCLES To FAILURE12 Effect of homogemzation on fatigue strength.
3 81 mm dia. centre specimens
homogeni/mg at 1315 C for I3h. Therefore, the notch efTecl of the surface cavities would override any impro\emenI resulting from modifications in dendnle morpholog) or more uniform distribution of segregated elements
In contrast, edge specimens treated similarly showed a marked improvement in fatigue strength (Fig.13). The fatigue strength was increased to the same level as that of centre specimens after HIP treatment (Tables 5 and 6|.
Some investigators' * M have shown that high- temperature homogenizjlion has little effect on the yield strength and L TS. but does improve ductility and low- lemperature notch toughness Therefore, the improvement m fatigue strength resulting from homogenizahon cannot be attributed to an increase m UTS. An improvement in the fatigue properties of an alloy cast steel homogenized at 12X0 C for 50h has been reported. 1 * No reason was advanced for ihe increase, but it has been suggested 1 " that the improvement in mechanical properties after prolonged high-tempera ture homogenizing may be due to the spheroidization of sulphide non-metallic inclusions which is know n to occur The effect of homogemzation on the fatigue strength of cJsl steel is being investigated further If the beneficial effect reported here is confirmed, then the use of HI P treatment on cast components will result in an improvement in the fatigue strength, both in the centre and at the surface, due to simultaneous consolidation and homogemzation
expectci.
Table 6
since Basof micrupor
Effect of
aran a .osit\ was
homogen
/ ' : have shonl> slightl)
»«n that thereduced after
38CH 1
|360 5
zing and normalizing on j£fatigue limit of 3 81 mm dia spec mode, rotating bending, SmMn -
Positron
Edge Edge
CentfeCentre
Treaimeni
Normalised at Homogenizedand normalize-Normalised atHomogenized
920'Cai 1 100 C
d at 920 C
31 1 TOO C
Fatigue limit MN m 2
285 310
250230
imens, stress 0
Fatigue ratio (FL UTSI
042 046
0 37034
and no-malp/ed at 920 C
Ul
£ 320
T3300S 280
-
"
1105
13 Effect3 81 mn
_— ————————— j ————————————— ! ————————
\ o normalized at 92O°C \» * annealed at 11CO°C and -
\ normalized at 92O°C
0\
V "^-^_______^s^ ^^-^— .i i106 107
CYCLES TO FAILURE
*--»•V 2
o-«.10"
of homogenization on fatigue strength> dia edge specimens
Metals Technology September 1982 Vol9
354 Strode ei at Effect of pressing on fatigue strength
Conclusions References
1. The presence of surface microporosity has little effect on the yield stress and UTS of cast steels, but reduces ihe elongation. RA. and fatigue strength
2 Hot isoslatic pressing <HIP| is an effective way of eliminating microporosity in cast bteel
3. The fatigue strength of cast steel containing internal microporosity is improved considerably by HIP treatment, and may attain a value equal to that of a comparable wrought steel.
4. The improvement in mechanical properties resulting from HIP ts due to the combined effect of consolidation and homogenization
5 The additional cost of HIP treatment may be a deterrent to its wider use. but may be amply justified where the absolute reliability of high-integrity components is a primary requirement.
Acknowledgments
The authors wish to thank Dr J D Davies. Director. Polytechnic of Wales, for permission to carry out the work. and Dr T. J. Griffiths. Senior Lecturer. Department of Mechanical and Production Engineering, for helpful discussions Thev also thank HIP (Powder Metals) Ltd, Chesterfield, for carrying out the HIP treatment
1 'Designing wiih high strength steel castings'. Publication M46; 1965. New York. Climax Molybdenum Co.
2. H D HANFS D A SFIFERT and C R WATTS. 'Hot isOSUtlCpressing'. Publ No MC1C-77-34. 55-68; 1979, Columbus.Ohio. Baitetle Memorial Institute
3 c P MI FLLtR and J R HI MPHRFY 'American mcials processingand fabrication techniques', led. R. M Silva), Slratford-upon-Avon. UK. March 1974. IRDCo Ltd. Newcastle upon Tyne.and Universal Technology Corp . Dayton, Ohio, Paper 6.
4. t a tVANS L i FBFRT and c w BRIGGS: Proc. ASTM, 1956, 56,979 1010
5 s 7 LRAM Trans.AFS. 1960. 68, 347-3606 w j JACKSON Br Foundr\man. 1957.50, 211-2197 s J WALKtR Foundry Trade J., 1969. 127, 943-9508 P c FORREST 'Fatigue of metals', 135-146. 1962, Oxford,
Pergamon Press9 s b FROST K I MARSH andL P POOK : 'Metal fatigue'. 136-149;
1974. Oxford University Press. 0 i STRODF unpublished work1 G F w\siFiFWSki and N R LiNBLAD 'Elimination of casung
defects using HIP superalloys - processing'. MC1C Report 72, Metals and Ceramics Information Center. Battelle Memorial Institute. Columbus. Ohio. 1972
2 V1 BASARAN T Z KATTAMS R MtHRABlAN and M C FLEMMINGS: Metall. Trans . 1973. 4, 2429-2434
3 p j *HEARN and F C OL'lGLEY J. Iron Sleel Inst.. 1966. 204. 16-22
4 j o kLR*and P c ROSENTHAL: Trans. AFA. 1946.54, 154-183.5. i- c OLIGLEY and P j *HFARS Mod Cast-, 1965. 47, 8136 F c 01 IGLEV and F OLICV 'Proc 1st Army materials
technology conf (ed J J Burke el al.) Wentworth-by-the-Sea,N.H..Oct. 1972.339-374
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An evaluation by the DGM's Metallography Subcommittee of data and knowledge built up in the field of Interference Layer Microscopy clearly presented to enable the metallographer to apply the different techniques reliably and with con fidence The Atlas is in two parts• an introduction giving the essential theoretical basis and informa tion necessary for the practical application of IL techniques; and• an illustrated section with high-quality colour plates, giving characteristic examples from groups of materials to show the possibilities of IL metallography Coating conditions are given for all examples, so that the metallographer can reproduce the contrast without excessive theoretical knowledge. The Metals Society has agreed with the DGM to distribute the Alias in the UK and various other parts ot the world outside Germany A fully des criptive folder/order form is available from the Society
Metals Technology September 1982 Vol 9
VOL73 PART 10 OCT 1980
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PresidentM. I. Danischewsky, FIBFSenior Vice PresidentR. CarsweU, MIBFJunior Vice PresidentJ. L. Younger, MIBFHonorary Treasurerti. E. Williams, BSc, FIBFSecretaryG. A. Schofidd, MIBF, Dip.MTechnical AdviserB. F. Boultbee, T.Eng. MIBF
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ContentsFoundryman's World
]4ih Scottish Weekend Conference31st National Works VisitsTredomen Engineering visitMid Gloucestershire TechnicalCollege visitJohn William Foundries Ltd. visitB.S.C. Hamiltons visit
Diary of events
Papers
275
277
287
292
302
23rd/24th October 1980 31»l National Works VllltsWales and Monmouth
5th November 1980 A one-day seminar Towards 1984: Robots in the Foundry IndustryEurocrat Hold, Maidenhead.
309
Noise abatement in foundnes. I. Eddington & N. Eddington. Primary austenite dendrites in grey cast irons. M. Ghoreshy, M. Zehtab- Jahedi & V. Kondic. The surface integrity and mechanical properties of electrochemicaJly machined cast steel surfaces. I. Strode & M. B. Bassett.Approaches to improved job scheduling in foundries. J. T. Southall & T. D. Law. Developments' in zinc die casting technology. A. J. Wad & D. L. Cocks.A sodium silicate bonded self- hardening sand system using a fluoride hardener. P. L. Jain & P. K. Panda.Morphology of unidirectional solidification front and undercooling in eutectic Fe-C-Si grey cast irons. T. Owadano, K. Kjshitake, M. Fujii i "& K. Miyamoto.47th International foundry congress, Jerusalem. Synopses of papers.
The surface integrity and mechanical properties of electrochemically machined cast steel surfacesI. Strode, M.Phil., C.Eng., M.I.M.. C.G.I.A., M. B. Bassett, B.Sc. M.ScTech Ph D C.Eng., F.I.Mech.E., F.I.Prod.E. " ' "IntroductionIn approaching the problem of metal removal, the present day engineer is faced with a-variety of machining methods from which to choose. Of the many criteria which must be considered before a final choree is made, an important and frequently overlooked factor is the effect of the machining method on the surface properties of the machined stock.
In recent years much research has been done on the effect of a number of conventional machining processes on the surface integrity and properties of the generated surfaces 1 2 . Similarly, the effect of the electrodischarge machining of some materials has been reported^. How ever, the corresponding study of surfaces produced by electrochemical machining has been limited to a few high chromium alloys 4 .
Electrochemical machining is now being increasingly used for metal removal, particularly for alloys which are difficult to machine by conventional methods. However, its use has been largely restricted to wrought alloys and little research has been done on the behaviour of cast alloys.
Since cast alloys are less homogeneous than wrought alloys, inferior surfaces and properties may result after electrochemical machining, due to ^elective electrolytic attack. In view of this possibility, a major objective of the present work was to investigate the more important factors which are likely to affect the surface integrity and mechanical properties of selected electrochemically machined cast steels and compare them with wrought steels of similar composition. Particular attention was given to the effect of current density on the surface structure and the fatigue strength in view of its known sensitivity to variation in surface integrity.
Experimental procedureMate rial5
The initial investigation was confined to.a medium carbon cast and wrought steel ha\mg the following chemical composition.
Table I
250 mm long and 100 mm square. In order to minimise the effect of microporosity, the fatigue specimens were taken from longitudinal sections cut only from the peri pheral areas of the casting. The macrostructure of the cross-section of the casting. Fig. I, confirms that all the fatigue specimens will have a columnar gram structure.
Heat treatmentThe as-cast blocks were sawn into 25 mm square
longitudinal sections and turned to 20 mm diameter prior to heat treatment. Both cast and wrought bars were normalised by heating in an atmosphere of gaseous nitrogen at 920~C for two hours followed by air cooling to room temperature.
Electrochemical machiningFatigue specimens of the dimensions shown in Fig. 2
were prepared using a profile lathe for the initial form. The final 0-25 mm was removed by electrochemical machining using a Herbert-Anocut 150 machine and a I 2 Molar solution of sodium nitrate with a pH of 7-8 and an electrical conductivity (Ke) of 0-10 ohrrr 1 cm~", operating at a temperature of 38 °C. The tooling consisted of a static cell which is shown in Fig. 3, located in the work area of the Herbert-Anocut machine. The split copper electrodes and the final form of the fatigue speci men is shown in Fig. 4.
A minimum of ten fatigue specimens were machined in accordance with the conditions given in Table II. Speci mens 'C' were machined at the highest permissible current density.
Table II
Identity
CD
Electrolytetemp. "C
3.838
Inletpressiirbar
6565
Machining e nme
seconds
1450
Vola
1510
Current density amps/cm2
Initial
39-053070
Final
335025-1
Material
Cast steel
Composition %
C Si Mn S f N, Cr Mo Sn
037 032 066 0-023 OO23 026 027 0-09 OO19 048 026 085 0044 0-032 012 010 <0-OI <0-OI
Cu
0 16 0 15
The wrought sleel was supplied in the form of 20 mm diameter hot-rolled bars and the cast steel as cast blocks
Mechanical propertiesThe tensile strength and related properties were deter
mined in accordance with BS: 18:1971 and Charpy V-I. Strode is Principal Lecturer in Ihe Department of Mechanical no tcn specimens were prepared and tested as specifiedand Production Engineering. The Poly,echmc of Wales. mBS' 131-1972 Faligue tests were carried out using anM. B Basse,, ,s Head of .he De,»r.men, of Mechanical and Pro- • ' m whjch ^ ^^ |oad (j duction Engineering, The Polytechnic of Wales. (F.I349). n *•'J
281
The surf net litlegrily and mechanical properties
applied at one end of the specimen by means of an oscillating spindle and measured by means of a torsion dynamometer attached to the other end
The mechanically polished specimens, Fig. 2, were prepared by conventional machining in accordance with the prescribed procedure in BS: 3518: Part 2: 1962, finishing with 6 jim and I (im diamond paste in the final polished condition both before and after stress relieving under a vacuum of - lO^Torrata temperature of 600 C
Surface integrityThe surface finish in the circumferential direction was
determined using a Taylor-Hobson Talysurf machine Additionally, electrochemically machined surfaces were examined by means of optical microscopy using a Nomar- ski interference contrast objective and also by a scanning electron microscope.
ResultsThe 0-2% proof stress, ultimate tensile strength and
related properties, including Charpy V-notch values of the heat-treated (normalised) steels is given in Table III Table III
Material
Wrought slcelCasl sleet
02; PSA/.V'm2
40441 1
</T\SMNIrrf
723673
F-long °'0S6S So
23 221 6
Rof I" g
48 423-2
CharpyV-noichimpact J
30-28-3030-32-32
The fatigue limit after the respective surface treatment is given in Table IV, including the fatigue ratio (UTS/FL). The results are also presented graphically in Figs. 5 and n.
ofeleclrocnemically machined cast utel surfaces
Discussion of resellsThe surface finish produced by electrochemical machin
ing is inferior to that obtained by the mechanical polishing of fatigue specimens, Table V When machining at 15 volts, which is the safe maximum capability of the machine, (he current density varied from 39-0 to 33-5 A cm', due to the increasing gap size This is lower than thai generally used in electrochemical machining, which may vary from 50 to 150 A/cm 2 . However, the surface finish of the wrought steel is of the same order as that obtained by Mao et al>, using a 3 Molar sodium nitrate solution. The surface finish of the corresponding cast steel is slightly inferior to that of the wrought steel as would be expected from its greater heterogeneity. This is also reflected in the microstructure due to the differen tial electrolytic attack in the case of the cast steel. Figs. 9 and 10. The surface finish deteriorated markedly with a decreased current density, particularly in the case of the cast steel. This is again indicated in the microstructure. Figs. 11 and 12, where differential attack of the cast steel revealed the cast dendritic structure. Fig. 12.
Mechanical properties, such as tensile strength and hardness are not very surface sensitive and are therefore little affected by electrochemical machining 6 . However, since fatigue is predominantly a surface phenomenon, a reduction in the fatigue strength would be expected after electrochemical machining. This may be due to the removal of surface compressive stresses, an inferior sur face finish, or to preferential electrolytic attack. There fore, it is important to know the contribution of each of these factors to the total reduction in the fatigue strength.
All fatigue specimens, however carefully prepared, will have surface compressive stresses which will result in an
Identity
ABCD
Surface preparation
Mechanically polishedStress relievedECM— IS voltsECM — 10 volts
Wrought 1
fatigue In.W/V m*-
360310310220
neel
miFatigue ratio
050043043030
Cast Steel
Fatigue lir° 0 Reduction MM.m-
- 26514 14 25039 220
nitFatigue ratio
040—037033
/o Reduction
_—5-7170
Surface finish measurements in the circumferential direction are recorded in Table V
Surface preparation
Mechanically polished
hCM — 15V
ECM— 10V
Ra v-m
Wrought steel
OO02 OO03
1 00 1 05
200 300
Casl steel
0-002 0002
1 375 I 125
2 620 3 750
The microstructure of the wrought and cast steels after normalising at 920 C is shown in Figs. 7 and 8. After electrochemical machining, the surface structure of the fatigue specimens as revealed by the Nomarski objective is shown in Figs. 9-12 and selected SEM micrographs in Figs. 13-16.
increased fatigue strength- Therefore, the true fatigue strength of an alloy is obtained either by electropolishmg or by vacuum annealing prior to testing. However, since electropolishmg also attacks sulphide inclusions in steel, the vacuum annealing method was preferred in this case. A reduction of 14° 0 occurred in the fatigue limit of the wrought steel after vacuum annealing and a similar re duction occurred after electrochemical machining at the maximum density. Table IV, It is therefore clear that under the conditions of the test, the reduction in the fatigue limit was mainly due to the removal of surface compressive stresses. Evans et al 7 have shown that com pressive stresses exist in Nimonic alloys up 10 a depth of about 02 mm and that the fatigue life at a constant stress level of _,* 433 MN m : decreased from 10 7 cycles in the highly polished condition to a constant value after the removal of 0 10 mm by electrochemical machining in a I?",, sodium chloride solution. At the lower current density. Table IV, a reduction in the fatigue limit of 40% occurred. It is therefore evident that with decreasing current density, selective electrolytic attack also occurs. This is confirmed by the microstructural evidence, Figs,
THt surface integrity and mechanical properties
1 1 and 13, which shows a greater degree of pitting, from which fatigue cracks may develop.
When the corresponding results for the cast steel are compared, the reduction in the fatigue limit was con siderably lower, 6° 0 and I5° 0 respectively for the high and low current density, This is contrary to expectation in vie« of the greater heterogeneity and inferior surface finish of the cast steel, Table V. This therefore clearly indicates that the cast steel is considerably less affected by the surface condition than the wrought steel.
Evans et al 8 , have shown that cast steels are less notch sensitive than comparable wrought steels. They found that whilst the fatigue limit of highly polished cast speci mens was about 20% lower than that of comparable wrought steels, the notched fatigue limit (Kt - 2-2) was the same for both steels. Also, whilst the fatigue limit of highly polished and lathe turned cast steel specimens was of the same order, the fatigue limit of lathe turned wrought steel specimens was reduced by about 28" 0 when compared with the highly polished specimens, li is therefore evident that cast steels are less sensitive to the surface conditions than wrought steels and hence less
of etectrochemically machined ctm steel surfaces
0 =+0.25 for ECM specimens =+0.05 for mechanical polishing
yrRIZ.7
K 7.938p'925 /
* ——— ~ ——— *"• ————————————— r. ———————— -i
^t '
-<P^.JU
. 22 23 . _
50.8dimensions in mm
FIG. 2 DIMENSIONS OF FATIGUE SPECIMENS
FIG. CAST STEEL - MACROSTRUCTURE
FIG 3; STATIC CELL POSITIONED F
FIG. 4: SPLIT CWWR EUCTWOCS *D tCH FATIGUE SKCUCH
affected by electrochemical machining in spite of their greater heterogeneity.
Since the fatigue strength of steel generally increases as the tensile strength increases, realistic comparisons may only be made if the fatigue ratios are compared. Table IV indicates that the fatigue ratio in plane bending of 4 3 mm diameter highly polished cast steel specimens was 0-40 compared with a 0-50 for the corresponding wrought steel. These results are similar to that obtained by Evans et al 7 , determined on 5-6 mm specimens using R. R Moore rotating bending machines. This seems to suggest that differences in the fatigue limit of wrought steel due to the type of stress application and specimen size may not be applicable to cast steel due to its intrinsic characteristics. This is being further investigated.
When designing steel castings a lower fatigue ratio than that of comparable wrought steels must be used.
501H
20010! 10°
CYCLES TO FAILURE (Nl FIG 5 S-N CURVES FOR WROUGHT STEEL SPECIMENS
10' 10s 106CYCLES TO FAILURE (N)
FIG6 S-N CURVES FOR CAST STEEL SPECIMENS
101
FIG. 7: WROUGHT STEEL - NORMALISED x 200 FIG. 8: CAST STEEL - NORMALISED x 200
FIG. 9: WROUGHT STEEL. ECU 15V x 500 FIG. 10: CAST STEEL. ECU 15V x 500
284
r surfacr IKIegnil and mtdianital properties oft leclrocfle mi'-olli machine* ran Heel surfaces
FIG. 11: WROUGHT STEEL. ECM 10V x 500 FIG. 12: CAST STEEL. ECM 10V x 500
FIG. 13: WROUGHT STEEL. ECM 15V « $00 FIG. 14: CAST STEEL. ECM 15V x 500
FIG. 16: CAST STEEL. ECM 10V x SOO
rhe surface integrity and mechanical properties of etectrocnemicallv machined cast steel surfaces
However, the advantages of cast steel is that it is less affected by subsequent machining operations than wrought steel. This is also true of electrochemical machin ing, particularly at a high current density which only resulted in a reduction of 6° 0 in the fatigue ratio com pared with 14° 0 for"the corresponding wrought steel. When machining at a lower current density the fatigue ratio of the cast steel is higher than that of the corres ponding wrought steel which again emphasises the lower notch sensitivity of the cast steel. Further work is in progress to determine the extent of the notch effect pro duced by electrochemical machining.
ConclusionsIt has been shown that after electrochemical machining
using a 1-2 Molar sodium nitrate solution:1. An inferior surface finish is obtained, particularly in
the case of the cast steel.2. A reduction m the fatigue strength of both wrought
and cast steel occurred.3. Provided that the steel is machined at a suitable
current density, the reduction in the fatigue strength is mainly due to the removal of surface compressive stresses.
4. A reduction in the current density results in a further
deterioration in the surface finish and a substantial reduction in the fatigue strength due to selective electrolytic attack and excessive pitting of the surface In spile of its greater heterogeneity, the fatigue strength of cast steel is less affected by electro chemical machining than similar wrought steels due to their lower notch sensitivity.It is important when designing steel components that a fatigue ratio corresponding to the surface condition is used.
Field. M. and Koster, W., ASTME, Em 68-516, Jan. 1968.Symposium on Surface Integrity, PittsburghL. P. Tarsov and W E Liftman, ASTME. Em 68-517, Jan.1968. Symposium on Surface Integrity Pittsburgh1. A. Bucklow and M. Cole Met. Revrews, No. 135, Metalsand Materials, 3. 6. 1969, 103.G. Bellows and J. B. Kohls, Amcr. Soc.of Man. Eng., MRR —76-12, 1976, 1K W. Mao, M. A. La Boda and J P. Hoare, Jnl Electrochem.Soc., 119, 4, 1972, 419.J. A. Gurklis, Defence Maienals Information Centre, Report213. Battelle Memorial Institute. 1965J M tvans, P. J. Boden and A. A. Baker, Proc. 12th Int.Mach. Tools DCS. Conf., 1971, 27).E B. Evans, L J. Eberl and C W Bnggs, Proc ASTM. 56,1956, 979.
286
