Medical, Pharma, Engineering, Science, Technology and Business

University of Oldenburg, Germany

- *Corresponding Author:
- Gerd Kaupp

University of Oldenburg, Organic Chemistry I

PO Box 2503, D-26111 Oldenburg, Germany

**Tel:**+49 4486/8386

**E-mail:**gerd.kaupp@uni-oldenburg.de

**Received Date**: June 09, 2017; **Accepted Date:** June 17, 2017; **Published Date**: June 27, 2017

**Citation: **Kaupp G (2017) Challenge of Industrial High-load One-point Hardness
and of Depth Sensing Modulus. J Material Sci Eng 6: 348. doi: 10.4172/2169-0022.1000348

**Copyright:** © 2017 Kaupp G. This is an open-access article distributed under the
terms of the Creative Commons Attribution License, which permits unrestricted
use, distribution, and reproduction in any medium, provided the original author and
source are credited.

**Visit for more related articles at** Journal of Material Sciences & Engineering

The physics of industrial single-point force indentation hardness measurements (Vickers, Knoop, Brinell, Rockwell, Shore, Leeb, and others) is compared with the depth-sensing nano, micro, and macro instrumental hardness technique that provides several further mechanical parameters, when using the correct force/depth curves exponent 3/2 on the depth of the loading curves. Only the latter reveal phase change onset with transition energy, and temperaturedependent activation energy, which provides important information for applications of all types of solids, but is not considered in the ISO or ASTM standards. Furthermore, the high-load one-point techniques leave the inevitably even stronger and more diverse consecutive phase-transformations undetected, so that the properties of pristine materials are not obtained. But materials are mostly not (continuously) applied under so high load, which must lead to severe misinterpretations. The dilemma of ISO or ASTM standards violating the basic energy law, the dimensional law, and denying the occurrence of phase changes under load is demonstrated with the physics of depth-sensing indentations. Transformation of iterated ISO-hardness and finite element simulated hardness to physical hardness is exemplified. The one-point techniques remain important for industry, but they must be complemented by physical hardness with detection of the phase transformation onset sequences for the reliability of their results. The elastic modulus EISO from unloading curves as hitherto unduly called "Young's" modulus has nothing in common with unidirectional Young's modulus according to Hook's law, because the skew tip faces collect contributions from all crystal faces including shear moduli, while iteration fit is to Young's modulus of a standard. Unphysical and also physically corrected multidirectional indentation moduli mixtures of mostly anisotropic materials and there from deduced mechanical parameters have no physical basis and none of these should be used any more. A possible solution of this dilemma might be the use of indentation-Ephys and bulk moduli K from hydrostatic compression measurements. The reasons for obeying physical laws in the mechanics of materials are stressed.

Brinell hardness; Elastic modulus; Force-depth curves; Hook's law; ISO and ASTM standards; Macroindentation; Physical hardness and modulus; Rockwell hardness; Ultrasound; Vickers hardness, Young's modulus

While present depth-sensing indentation hardness and modulus
determination is obtained by nano- and sometimes micro-indentation
(nN to μN and mN), instrumented macro-indentation is also possible
up to 80N. Industrial non-depth-sensing techniques still concentrate to
longer known Vickers (HV), Knoop (HK), Brinell (HB), Rockwell (HR),
Shore, rebound LH (Leeb), or more specialized macro-hardness tests.
These measure the impression diagonal, or diameters, or final depths
under specified conditions although with subsets for certain types of
materials. One subset of HV is the UCI technique (ultrasonic contact
impedance, requiring elastic modulus E_{eff}) measuring vibration damping
of a swinging stick with a Vickers diamond at the end, as inserted at a
predefined load. All these single-point high-force techniques require
1:1-calibration with test plates of closely related materials of "known"
hardness, also for canceling out not specified tip end radii. Several
hand-hold devices exist, which is practical in the steel industry. Handhold
equipment includes UCI, Leeb, Rockwell clamp, Brinell clamp,
Brinell Poldi hammer, etc. All of these techniques use rather empirical
definition and ISO (International Standard Organization) or the now
compelling ASTM (American Society for Testing and Materials)
standards. Between HV, HK (low load range), HB, HRB, and HRC
exist approximate conversion equations. This indicates relationships
between them. The equipment software usually calculates and displays
most of the different types of macro hardness. However, conversion
between them is often not precise enough for construction purposes,
notwithstanding the sometimes large experimental uncertainties, due to often low reproducibility between different user sites. And there is
serious non-compliance with basic physics.

The comparison with depth-sensing instrumented indentation according to ISO 14577 where three major flaws occur in the universal, ISO, and finite element (FE) simulated hardness, is difficult. The instrumented depth-sensing could recently be corrected for providing the physical hardness, eqn. (1) (where k is the slope of the so called "Kaupp-plot" eqn. (2)) by removing three physical flaws inherent to ISO 14577 [1-3].

H_{phys}=0.8k/π(tanα)^{2} (1)

F_{N}=kh^{3/2} (2)

Corresponding violations of physical laws have not yet been considered in the single-point-load techniques, but these must equally exist. This bears an important risk for the mechanics quality of industrial goods. A prevailing source of uncertainty is the non-considered phase transformation of materials that change the material's hardness and other mechanical properties, under the very large local pressure. It is well-known [1,3] that phase changes occur already at nanoindentation and lower micro-indentation. They must therefore be even more common in the macro range. Furthermore, the possibilities for detection of hidden horizontal cracks (except when these occur upon unloading) are not evident. All of the industrial indentation techniques also penetrate vertically onto flat surfaces, but now with a defined holding time at the predetermined force. Creep is assumed to be negligible. Some of these macro-techniques (HV, HK, HB) measure diagonals or diameters of the impressions that is left at the surface, others (HR, Shore) the indentation depth. But final depths can also be calculated from the indenter geometries in the former cases. The physical flaws as detected in the instrumented depth-sensing should be the same in all macro-hardness tests. Clearly, the depth relates to the diagonal or diameter left at the surface. So there is no principal difference to more precise depth sensing, except that the applied forces are usually very much higher. Very detailed and constantly refined ISO and ASTM standards are available. A comparison between these and the depth-sensing techniques is thus in urgent order, by applying the physical news from the nano- and micro-indentations [3].

Similar difficulties with elastic moduli concern only the depth
sensing unloading. The same dimensional energetic and phase change
violations of ISO standards can be principally corrected. However, it
turns out, there is not the claimed correspondence of ISO or physical
indentation moduli with Hook's Young's moduli, so that E_{ISO} should
no longer be iterated (Oliver-Pharr method), falsely called "Young's"
modulus, and used. Even E_{phys} is only a counterpart of H_{phys}, the physical
hardness. It will however be suggested to use bulk moduli instead.

The nanoindentations onto a polished optical disc 2 mm thick
NaCl single crystal (purchased from Alpha Aesar GmbH Co KG,
Karlsruhe, Germany) were performed at a Triboindenter^{(R)} with AFM
of Hysitron Inc, Minneapolis, USA, with proper calibration at 23,
100, 300, and 400°C (average of eight measurements). The author's
nanoindentations used a fully calibrated Hysitron Inc. Triboscope^{(R)} instrument with AFM in force controlled mode also with a Berkovich
diamond (R=110 nm). The cited literature data have been carefully
searched and interpreted in view of the generally deduced physical
laws, in accordance with validated experimental data. Phase changes
under load are detected by kink-type discontinuity [4] in so-called
"Kaupp-plots" according to eqn. (2) [1-5]. The precise intersection
point is obtained by equating the regression lines before and after the
onset of the phase change. The regression coefficients are calculated
with all 400-500 or 3000 original data point pairs using Excel^{(R)}, but excluding those from initial surface effects. Digitizing 50-70 almost
uniformly arranged data pairs were obtained from published loading
curves with the aid of the Plot Digitizer 2.5.1 program (www.Softpedia.com). The distinction of experimental from FE-simulated loading
curves succeeded with the "Kaupp-plot". The necessary correction to
comply with the energy law [1,3] is by multiplying the slope k (mN/
μm^{3/2}) with 0.8.

**Dilemma between ISO standards and physics**

The comparison of depth-sensing instrumented ISO-hardness with non-depth-sensing single point high-load techniques reveals undeniable physical similarities. The industrially used macro indentation techniques are governed by the same physical laws as depth-sensing nano to macro indentations. Unfortunately, present ISO standards are at variance with the corresponding physical laws [1,3] and the possible corrections of previously published indentation data require a detailed discussion here. The physical requirements for singlepoint load indentations reveal equally from the precisely determined facts of the better controlled depth-sensing continuous indentations, including the macro-indentation ones.

**Table 1** compares the depth-sensing hardness values of H_{phys},
H_{ISO}, and H_{simulated}, to demonstrate the importance of correct depthsensing
evaluation. It is also shown how the latter two can be corrected,
provided that the loading curves were published as for example in ref.
[6]. This is a practical application for the conversions of FE-simulated
or ISO hardness values (energy law violations and incorrect exponents)
into physical hardness.

Entry | Technique | h_{max}^{n} |
k or h_{max}Fmax(a) |
Hardness calculations and corrections |
---|---|---|---|---|

1 | Experimental curve linear regression | h_{max}^{3/2} |
k=5.9540(mN/µm^{3}/2) (energy corrected)(b) |
H_{phys}=k/𝜋tan∝2=0.24295(mN/µm^{3}/2) Independent on FN and h_{max} (no phase trans.) |

2 | Iterated H_{ISO} with 2/3 factor |
h_{max}^{2} |
- - | H_{ISO}=0.716 (GPa) × (2/3)≈0.477 (mN/µm^{2}) (still unphysical dimension, h_{max} unknown) |

3 | FE-simulated not corrected | h_{max}^{2} |
h_{max}=0.250 µm F _{max}=0.912 mN |
H_{simul} (as H_{univ})=0.6016 (mN/µm^{2}) |

4 | FE-simul. h_{max}^{1/2}no energetic corr. |
h_{max}^{2} |
h_{max}=0.250 µm F _{max}=0.912 mN |
H_{simul} (as H_{univ})=FNmax/𝜋tan 𝛼2 h_{max}3/2=0.2977 (mN/µm^{3}/2) (still energy law violation!) |

5 | FE-simul. 2/3; no exponent corr. | h_{max}^{2} |
h_{max}=0.250µm |
H_{simul}-corr2=2x 0.6016/3=0.4011 (mN/µm^{2}) (wrong exponent) |

6 | FE-simul., h_{max}^{1/2} and energetic corr. |
h_{max}^{2} |
h_{max}=0.250 µm F _{max}=0.912 mN |
H_{simul}-phys=0.8×0,2977=0.2382 (mN/µm^{3}/2) |

**Table 1: **Comparison and correction of *H*phys, *H*ISO, and FE-simulated H_{simul} loading curves ofAl [6] including the corrections in accordance with the exponential and energy laws; extendedtable from ref. [3]. (a)simulated parameters are not italicized; (b)energy correction factor 0.8.

Entry 1 shows the correct value Hphys, according to the Kaupp-plot with linear regression from the experimental loading curve in ref. [6].

Entry 2 deals with the published iterated HISO value that enormously differs in value and dimension. The difference is still very large when the energetic law violation is removed (based on the falsely believed "h2" the energy or force loss for the indention calculates to 33.33% energy law violation) (cf. refs. [1,2]). This is incomplete correction. It is not clear, which FN/h pair was used in the HISO iteration. Complete correction would also suffer from the exhaustive iterations that cannot be reverted.

Entry 3 deals with the FE-simulated hardness without the necessary
corrections: again, a large deviation in value and dimension from H_{phys}.

Entry 4 demonstrates only the dimensional correction, as h_{max}^{2} was
used instead of h_{max}^{3/2} for the relation with Fmax, but it is clearly not sufficient.

Entry 5 similarly reveals that only the removal of the energy law
violation (for believed h^{2} only 2/3 of F_{N} is available) is not sufficient.

Entry 6 shows that only both corrections (h^{1/2} and thereafter 0.8)
give a good value. But it has to be checked if all FE-simulations with
different iterations will give equally good correspondence with the
correct experimental H_{phys} value. All of the corrections in **Table 1** are
equally valid for the one-point non-depth-sensing macro techniques.
However, there is another very important flaw: phase changes under
load. Their occurrence and onset load can only be detected by depthsensing
indentation and Kaupp-plot.

As hardness is for the first time a physical quantity, there is no
possibility to change the physical hardness dimension say by division
with h^{1/2}, which would mean to accept the dimensional violation. The
inherent dimension has its meaning for all applications of the k-value
(force/depth^{3/2}, mN/μm^{3/2}) that cannot be dismissed. These include
depth-force relation (mar, wear, tribology), physical deduction with
elementary mathematic, adhesion work, pull-off curves, safe ratings
of materials, correlation coefficients with >3 nines or less noisy
with >4 nines, quantitative far-reaching energy, phase transition
onset, indentation work, compatibilities, transformation energies,
activation energies, creep, size effects, maximal load for reasonable
unloading curves, initial surface effects, high sensitivity by linear
regression, tip normalization, tip rounding effects, materials gradients,
inhomogeneous materials, geodes, crystal defects, edge interface, too
close impressions, grain boundaries, cracks, alternating, improvement
of FE-simulations, avoidance of polynomial iterations or varying
broken exponents, correct visco-elastic-plastic parameters, nanopores,
micro-voids, alternating layers, blunted tip effects, correction of false
mechanical parameters that rely on h^{2}, elementary mathematics instead
of iterative fitting, avoiding violation of the basic energy law with factor
0.8, failure risks with false mechanical parameters, tilted impressions,
faulty standards with phase change, all types of solid materials and plasticization types [1-5,8], quantitative sound mathematical basis,
universal validity, distinction of FE-simulated from experimental
loading curves, no denial of phase changes, and daily risk with
unphysical mechanic parameters [3].

**Macroscopic depth-sensing, hardness and phase transitions**

Depth-sensing nanoindentation extends up to 10 mN load, microindentations
up to 1 N, mostly with Berkovich indenter. Extensions to
macro- indentations have been achieved with Vickers up to 80 N for
soda-lime glass [7], but rarely repeated. Eqn. (2) has been experimentally
secured and physically deduced. It secures Kaupp plots for all of the
force ranges [2,5]. The dilemma of physics and ISO 14577, still believing
in "h^{2}", is clearly evident from **Figure 1**. It indicates exclusion of h^{2} and
phase transitions of soda-lime-glass, sapphire, and sodium chloride in
the macro-indentation range (the nano-ranges are "hidden" at these
macro-indentation ranges). The kink-type discontinuities (the phase
transition onsets) are at 15.37, 10.43, and 2.469 N load, respectively.

**Figure 1:** Kaupp-plot (FN versus h3/2) of published depth-sensing macro-indentations [5] showing kink-point intersections (phase transformation onsets) and linearity up to 80 N load; an "h2 relation" as believed in ref. [7] is excluded; the data have been extracted from ref. [7]; the correlation coefficients for all six regression lines are r>0.999.

The most important advantage of macro-depth-sensing is the detection of secondary phase transitions at very large forces. In the case of NaCl, the also endothermic fcc to bcc transition at 4.233 mN [8] is hidden at that scale. The same is true for a transition onset of sapphire at 26.5 mN load and for soda-lime-glass at 4.81 mN [5]. The other macroindentation techniques have the disadvantage that they cannot detect phase transition onsets: they practically always measure the hardness from (often secondarily) phase transformed materials, as embedded in the original material. They do not characterize the pristine materials! Clearly, one needs nano, micro and macro depth-sensing indentation in addition to the technical ones, for judging the materials mechanics. Most of the time, materials are not under such very high pressure. Also, the primary transition onset is important for material's mechanics (e.g. failure or fatigue).

The comparison of hardness measurements of sodium chloride
is particularly revealing, because (as in the case of sapphire and
soda-lime-glass) two consecutive phase transitions are involved. The literature knows Vickers microhardness data from the list for NaCl
properties of the MaTecK-Material-Technologie and Kristalle GmbH
collection (Jülich, Germany), reporting 0.20 GPa. Probably, this is the
same value as cited [7], but F_{max}/h_{max} is not known. HISO=0.252 GPa
(±2%) was recently measured at F_{max}=10 mN, but uncorrected after
its fcc-bcc phase transition (from the loading data for ref. [8]). After
energy and exponent correction before the phase transition onset this
gives with 0.8 k=5.8229 the H_{phys} value of 0.2376 mN/μm^{3/2} (GPa μm^{1/2})
(not violating the energy law etc.), as calculated with eqn. (1) [3]. H_{phys} is only obtained by linear regression of original data pairs without
any of the iterations for HISO. The nanoindentation up to 10 mN load
(sharp Berkovich, 1.17 μm depth at 10 mN force) creates the halite to
cesium chloride type phase transition (fcc to bcc) with onset at 0.697
μm and 4.233 mN load. It requires +0.04418 μJ/μN phase transition
work [8]. Ref. [8] reports also the activation energy (23 to 400°C) of
this first transition. The preferred hydrostatic transition pressure is
known as 26.8 GPa [9]. The calculated second phase transition (bcc to
layered CrB-type NaCl space group Cmcm) is hydrostatically expected
at 322 GPa, metallic from 584 GPa [10]. Most probably, the second
transition corresponds with the kink in the Kaupp-plot at 2.49 N load
and 21.1 μm depth [8] according to the loading curve of ref. [7]. The
transition work is +3.647 μJ/μN, which is very large when compared to
+0.04418 for fcc to bcc of NaCl, or for example +0.007066 for SrTiO_{3},
or −0.01126 μJ/μN for -SiO^{2} [8]. The discontinuity at 21.31 μm depth
of a sharp Vickers is a candidate for the predicted Cmcm phase of NaCl.

**Vickers hardness test and other one-point-load macrohardness
tests**

The load for HV varies in three ranges from 0.1 N up to 1500 N
(HBW10/3000 even with 30000 N; the W indicates tungsten-carbide);
the normal range is 40 - 980 N (HV4 - HV98). The Vickers hardness
test is most similar to the pyramidal instrumental depth-sensing, as the
Vickers indenter can be used in both techniques. One indents to the
chosen load, holds for 10-15 s (now 14 s), unloads, and calculates HV
from the average of the diagonals d of the impression. The standard
is given by eqn. (3), where FN is in kpf (kilopond force, it is a very old
standard) and d the diagonal length in mm of the residual impression.
Then, after conversion to N/mm^{2} units one reports (m HVn), where m
is the hardness value and n the vertical load F_{N} in kpf.

HV=1.86 F_{N}/d^{2} (3)

The first flaw deals with the dimensional error. Since the depth h
is geometrically related with HV's impression diagonal d according to
eqn. (4), the d2 relates again with h^{2} rather than with h^{3/2}. This gives
again a faulty inherent "FN∝ h2" relation (instead of the physically
deduced eqn. (2)), as in the instrumented depth-sensing force-depth
curve [2,3].

(4)

Next to the dimensional violation there is the second flaw:
violation of the basic energy law. The applied load is not only used
for the indentation depth but with 20% force and thus energy loss
(physically correct h^{3/2}): the sum of stress formation and plasticization,
including sink-in or pile-up, requires energy (if correction of h^{2} to h^{3/2} is not performed, the energy loss would be 33%) [1,3]. Long-range
features, often with pile-up around the square impressions, have long
been seen. Their universally quantitative occurrence in addition to
the created stress (Wlongrange) derives from the physically deduced ratio
of the different work contributions in eqn. (5) [1]. Clearly, the nonconsideration
of pressurizing and plasticization work is violation of the most basic energy law! The same is true with HB, HR, Shore, rebound,
and the techniques that use spherical impression instead of pyramidal/
conical ones, because these must also obey the physical relation of
eqn. (2) (h^{3/2} instead of h^{2}) as quantitatively deduced for depth-sensing
indentation. Also the UCI-Vickers hardness values, using ultrasound
frequency, suffer from the same flaws.

W_{applied}/W_{indent}/W_{longrange}=5/4/1 (5)

The third flaw is even more severe than in depth-sensing indentation,
because the forces/works and depths are much larger (compare the
NaCl, sapphire, and soda-lime-glass cases in **Figure 1**. Inevitably, there
must be several endothermic or exothermic phase changes following
each other, not to speak of hidden horizontal cracks that can also
occur upon pressure release at the unloading. Furthermore, one-point
measurements (rather than linear regression of loading curves with
Kaupp-plot) bare the risk of uncontrolled errors. This fact makes it
difficult to judge the reliability of HV etc. measurements that could in
principle be corrected for energy law and dimension (requiring depth
with tip rounding correction), but not with respect to force dependent
phase transformations under pressure, the detection of which require
analyzed force/depth curves with the physically founded exponent 3/2
on the depth (eqn. 2) [2,5].

The interpretation difficulties are demonstrated for HV
measurements with the test material 316L stainless steel. The general
claim is that HV values must not depend on the load. A publication
of 2016 gives a value of 281.6 HV0.1 (N/mm^{2}) at 0.981 N load [11].
The rounding of the Vickers pyramid is not given (its influence is
eliminated by comparison with test-plates impressions), but we
calculate for ideal Vickers a depth of 3.66 μm. Another publication of
2016 reports 280 HV^{3} (N/mm^{2}) at 29.43 N load [12] at a calculated
depth of 20.12 μm. Important questions are: why is the value for the
much deeper and 30 times higher force smaller by 1.6 MPa? Could
it be experimental error (this is calibration at a test material!), or
was the tip rounding too different, or are there undetected cracks,
or are consecutive phase transitions at the 30 times higher force
exothermic? Such considerations are missing, but the 208 HV3 value
was also converted into 217 HB; 95 HRBmax, and 89 HRB. Numerous
calibration tables exist and equipment software often displays such
converted results as well. The most important of the conversion
formulas that interconnect the various techniques are listed in **Table 2**.
Such conversions are termed "approximate" (conversion norm: ± 3.5%
of HV), but their use indicates their correlation. That means: all of the
single-point macro-indentations exhibit the same flaws with respect
to physics, notwithstanding the apparent technical problems. Clearly,
size effects due to phase changes are assumed to stay within the large
error allowance, and the end radii of the Vickers and Knoop pyramids
are not taken particular care of. Apparently, **Table 2** is only valid for
the same force, and these techniques are by no means universal, but
they need for every material a separate test sample with "known" HV
that must have been agreed upon. The phase change events are not
considered and neither can they be detected by the 1:1 calibration, even
though the forces vary from 0.1 to 150 kpf. Conversely, depth-sensing
is universally applicable to all solid materials but requires knowledge of the tip rounding that should be small enough, so that its influence can
be treated and corrected for as initial effect. ISO uses iterative relation
to a standard like fused quartz or aluminum. However, the physically
founded depth-sensing obtains absolute hardness values without test
samples [2,3], and it detects phase transitions directly as in **Figure 1** and [5].

HV to HB |
HV≈1.05 HB |

HV to HK | HV≈HK (low load region) |

HB to HV | HB≈0.95 HV |

HRB to HB | HRB≈176-1165/HB1/2 |

HRC to HV | HRC≈116-1500/HV1/2 |

**Table 2:** Some conversion formulas for one-point-load hardness values.

These considerations clearly indicate the close relation of the empirical single-point load macro-techniques that use either surface (HV, HK, HB) or depth measurements (HR, Shore, etc.) to the instrumental nano- and micro- and macro-indentations. Thus the same flaws, as in the ISO standards or FE-simulations, as based on the Oliver-Pharr technique, are involved: first the violation of the basic energy law, second the wrong dimensional error (violation of the Equation 2), and third the non-consideration of phase transformations that cannot be detected during the load and hold periods. The high load capabilities of depth sensing should be extended above 80 N. Clearly, depth-sensing measurements should always be separately available for every material's charge in addition to the fast HV, HR, etc. measurements for rapid and on-site production control, in order to avoid risks from unrespected phase change onsets giving polymorphs with different mechanical properties. And the study of crack onsets is important.

**Tension, compression and speed of ultrasound for Young's
modulus E**

Elastic moduli cannot be obtained by one-point-load hardness tests,
but ISO iterates it with depth-sensing unloading. The transformation of
E_{r-ISO} into "Young's" EISO from exhaustively iterated unloading curves is
achieved with eqn. (6), where both the Poisson's ratio and modulus of
the material and the indenter (diamond) occur. This gives values with
unchanged dimension but still burdened with the violating of physical
laws by the three major physical flaws (dimensional, energetics, unclear
solid phase). ISO calls such E values from unloading curves "Young's"
moduli.

There may however be severe objections against equating
indentation moduli E_{ISO} with Hook's Young's moduli. This holds also
for the indentation E_{r-phys} and with eqn. (6) E_{phys} (the pendant to Hphys)
[3] with different dimension (GPa μm^{1/2}) [7].

(6)

It does not help that the UCI-Vickers hardness test uses ultrasound
response, which requires an effective elastic modulus E_{eff} from
calibration tables for consideration of the E-module. UCI is not a
technique for modulus measurement. The reason for eqn. (7) is the
universal eqns. (2) and (5) for indentations, which means long-range
work for pressurizing and plasticization consumes 20% of the applied
work, and thus force, in case of correct dimension according to eqn.
(2) (or 33% as long as the false exponent 2 on h would be applied). But
the use of E_{phys} requires some efforts with the calculation of the initial
slope of the unloading curve using the original data, rather than a ruler
to the recorded curve.

(7)

In the absence of original data it can appear impossible to
graphically approach the initial slope ΔF_{max}/Δh_{max} that is the iteration
result by Oliver-Pharr. It the claimed EISO value of 73 GPa (Berkovich,
R=50 nm) from ref. [6] up to 215.8 nm followed by creep up to 266
nm depth. Actually, ISO iterates A (projected contact area) with an
unrelated standard for final height hf=h_{max}- 2F_{max}/S for A and fits 80% or 50% of the exponential unloading curve iteratively with FN=A(h_{max} - h_{final})^{m}, where A, h_{final}, and exponent m (between 1 and 3) are the
free parameters. Stiffness S at peak load is then obtained by the
differentiation dFN/dh=S=Am(h_{max}-h_{final})^{(m-1)} for obtaining the maximal
slope. This circumvents the slope detection. E_{r-ISO} is then calculated as π^{1/2}S/2 A_{hc}^{1/2} and the result is called "Young's modulus" after application
of eqn. (6). This is objectionable ISO standard.

The principal problem with such definition of an indentation
modulus is the anisotropy of most materials that cannot be tackled by
indentation, irrespective of the possible physical corrections eqn. (7).
For example, it is known from the fact that different faces of a crystal
give different E_{r-ISO} moduli depending on the different predominance
of the crystal faces towards the tip (e.g. αr-SiO_{2} varies E_{r-ISO} between
105.0 and 133.6 GPa onto 5 different faces) [13].The skew indenter
surfaces collect in fact a mixture of some sort of different elastic
moduli from all of the different directions around the tip and there
are also shear-moduli involved upon the unloading. This is far away
from unidirectional Young's modulus, depending on Hook's law
eqns. (8) and (9). Thus, EISO is incompatible with Hook's law, and
indentation-E_{phys} can also not be made compatible. Any similarities of
EISO values with Young's moduli are thus fortuitous. They derive from
the iterative fitting to the unidirectional Hook's value of a standard.
They are therefore fortuitous, because of both the multi-directionality
and because of the striking physical errors of E_{ISO}. They do not have
the same meaning, as might be suggested by the unfortunate common
wording. Fortunately, an extensive amount of well-studied Hook's
Young's moduli for all independent directions of preferably cubic and
other high symmetry crystals are tabulated and do not need repetition
by indentation. The complexity of the 6x6-matrix treatment of Young's
moduli, leading by some matrix symmetry to generally 21 independent
elastic constants that are further reduced by crystal-symmetry to 9, 7, 6
and in the cubic case 3 independent moduli has been amply described
(for example in ref. [14]). So it is suggested to call E_{phys} eqns. (6) and
(7) "indentation modulus" and check, whether the three-dimensional
bulk modulus, as obtainable from hydrostatic pressurizing, is an equal
or superior parameter for characterizing the elastic properties of micro
or macro materials.

It is essential now to briefly repeat the Hook's technique for obtaining Young's moduli E, where the shear modulus detection is excluded. The clearest experimental determinations of E are by tension/ compression eqn. (8) or ultrasound speed eqn. (9). The uniaxial tension or compression gives the simple elongation/depression Hook's law eqn. (8), as long as these are fully reversible. Transversal thinning/ thickening is always mentioned, but transversal work can apparently be neglected. L is length, p is the generated pressure (force per area), E Young's modulus.

(4)

Eqn. (9) recalls the ultrasound speed technique in long rods with
diameters smaller than the ultrasound-wavelength, excluding shearwaves,
where frictional loss may be small or ineffective. It is used for
the longitudinal speed *v*_{s} in such rods, where E is Young's modulus
and ρ is density. These and more complicated Hook's techniques are
generally accepted textbook physics.

(9)

For practical reasons we regret that the Hook's law techniques require much larger test samples with highly specialized geometric shape. They are therefore more difficult to perform and less versatile than would be indentations, that appear however inappropriate for E. The present situation is at best exemplified with the simplest case, cubic isotropic aluminium.

We have to distinguish tabulated Young's modulus (E=69 GPa),
shear modulus (G=25.5 GPa) and bulk modulus at hydrostatic
compression (K=76 GPa). This compares to claimed invalid EISO=73
GPa [6] that must be decreased to 10.7 GPa by making the physical
corrections. Clearly, nothing from the unloading is fitting with the
reliable Hook's values. There is no hope left that indent-E_{phys} (mN/
μm^{3/2}) values could be converted into Hook-E_{phys} (mN/μm^{2}) values
(for example by division with h_{max}^{1/2}), because they would have totally
different meaning. Again, it does not help that EISO is iteratively fitted
with respect to a unidirectional Hooks's Youngs modulus of a test
material.

The consequences for the recent use of physically unsound E_{ISO} values are detrimental, when their use for mechanical parameters is
considered. The particular importance of an indentation modulus is
evident from numerous applications. The listing 1 through 12 indicates
various examples.

1. All elastic properties

2. Input for FE-simulations

3. Stress-strain response

4. Film hardness and film adhesive strength

5. Strain hardening

6. Creep calculation

7. Material fatigue, fatigue strength

8. Adhesion calculation (DMT or JKR)

9. Viscoelasticity studies

10. Sliding friction coefficient

11. Contact area at dynamic testing in continuous stiffness mode

12. Fracture toughness

At present it appears only possible to calculate Young's modulus
E of new materials for certain directions and test the quality of such
calculations with as close as possible materials, for which the Hook's
values are known, or to rely on indentation-E_{phys} or on bulk modulus
K by hydrostatic pressure experiments for the consideration of reliable
elastic materials properties.

**Reasons for obeying physical laws**

It is very clear that mechanical properties must not violate basic physics, be it in academia, industry, medicine, or daily life. That does not mean that purely empiric methods like the Mohs hardness scale (who scratches whom) are also useful. However unphysical parameters must not violate physics. And one must not try to make physical correlations with unphysical parameters. For example, Mohs says steel cuts leather. However, there is also mechanochemistry that explains why barbers can sharpen their blades with leather [15]. Clearly, also the size of the components and the chemical composition of the solids play an important role (here polymers are tribomaterials) [15]. Brittleness, ductility, lubrication are further qualities apart from hardness and elasticity, that have their meaning in particular applications. Hardness and elastic moduli should be physical rather than empirical due to countless technical constructions where different materials must work together and alloys or composites must be compatible rather than fail upon short use. Materials are often used under low pressure where they are not phase-transformed. And different materials have their phase change onsets at varied pressures. This provides severe risks when they are perhaps only compatible under very high pressure as high pressure polymorphs, but not at lower or ambient pressure where they are at rest. Everyone knows that virtually all purchased goods with granted guarantee periods fail (shortly) after that period, or airliners must have very short control and replacement terms of all parts, because they must not fail. Only physically sound parameters of hardness and modulus with all of the numerous other mechanical parameters that depend on them should be used, instead of violating basic physical laws with HISO and EISO. The dilemma of ISO-standards against physics is a thread for daily life, because falsely calculated materials bear enormous risks for lifetime and failure. Some examples are composite materials (also solders) that may not properly fit together, or exploding turbines, or breaking windmill blades, or micro-cracks in airplanes and huge pressure vessels of power plants, or breaking medicinal bone implants due to incompatibility, etc.

The comparison of single-point load macro-indentations with
physical and mathematical precisely handled depth-sensing nano,
micro, and macro indentations reveals three major flaws of the former
that can be and have been removed for the latter [3]. All depends on
the physically deduced exponential law F_{N}∝h^{3/2}, instead of the believed
h2 from Sneddon, Oliver-Pharr, and ISO standards [3]. The same flaws
(violation of the basic energy law, dimensional error against physics,
and disregard of phase changes under load) are also inherent in present
ISO and ASTM standards that still do not apply basic physics from
the depth-sensing techniques. Since the one-point force techniques are
much more rapid and comfortable in industry, these purely empiric
techniques with standardized calibration necessities at test plates and
tables for different material types are now only acceptable, when the
materials in question have also been studied on the genuine physical
basis with force/depth curves, as described here and in ref. [3]. Depthsensing
ISO-standards are subject to urgent changes for complying with
physics. Most serious in view of failure risks are the present disregard of
phase transition (phase change) onsets, and size depending very large
differences between faulty HISO and the much more precise Hphys values
with different dimensions. Similarly, indentation elastic modulus E_{ISO} (falsely called "Young's modulus") fails: it suffers from the same physical
flaws and has no relation to unidirectional Hook's law. The unloading
skew pyramid or cone surfaces collect a mixture of multidirectional
elastic moduli and shear moduli. Therefore, indentation-moduli have
a totally different meaning than Hook's Young's modulus. They cannot
be given the same name, and the term EISO is also worthless due to
three physical flaws, and to questionable iterating fitting techniques as
initiated by Oliver-Pharr and taken up by ISO. The incredible claim that
ISO would deal with unidirectional Young's modulus has to be rejected.
It is not at all available for indentation unloading. E_{ISO} and deductions
there from are unphysical and their use must be discontinued. The
use of indentation-E_{phys} or bulk moduli K should be used in situations
where the one or the other appears more appropriate or better both
for the mechanical characterization of materials. Phase changes under
pressure must be controlled as detected from the mathematical analysis
of instrumented loading curves, so that the rapid single point high-load
indentations can find the appropriate interpretation.

We thank Dr. U. Hangen Aachen, Germany from Hysitron Inc for the indentions onto NaCl with a Triboindenter(R).

- Kaupp G (2013) Penetration Resistance: a new approach to the energetics of indentations. Scanning 35: 392-401.
- Kaupp G (2016) The physical foundation of
*F*N=*kh*3/2 for conical/pyramidal indentation loading curves. Scanning 38: 177-179. - Kaupp G (2017) The ISO standard 14577 for mechanics violates the first energy law and denies physical dimensions. J Mater Sci Eng 6: 321-328.
- Kaupp G, Naimi-Jamal MR (2010) The exponent 3/2 at pyramidal nanoindentations. Scanning 32: 265-281.
- Kaupp G, Naimi-Jamal MR (2013) Penetration resistance and penetrability in pyramidal (nano) indentations. Scanning 35: 88-111.
- Soare S, Bull SJ, Oila A, O'Neill AG, Wright NG, et al. (2005) Obtaining mechanical parameters for metallization stress sensor design using nanoindentation. Int J Mater Res 96: 1262-1266.
- Thurn J, Morris DJ, Cook RF (2002) Depth-sensing indentation at macroscopic dimensions. J Mater Res 17: 2679-2690.
- Kaupp G (2014) Activation energy of the low-load NaCl transition from nanoindentation loading curves. Scanning 36: 582-589
- Li X, Jeanloz R (1987) Measurement of the B1-B2 transition pressure in NaCl at high temperatures. Phys Rev B 36: 474-479.
- Chen X, Ma Y (2012) High pressure structures and metallization of sodium chloride. Europhys Lett 100 : 25005-1 - 25005-4.
- Wang D, Song C, Yang Y, Bai Y (2016) Investigation of crystal growth mechanism during selective laser melting and mechanical property characterization of 316L stainless steel parts. Mater Design 100: 291-299.
- Maestracci R, Sova A, Jeandin M, Malhaire JM, Movchan I, et al. (2016) Deposition of composite coatings by cold spray using stainless steel 316L, copper and tribaloy T-700 powder mixtures. Surface and Coatings Technol 287: 1-8.
- Naimi-Jamal MR, Kaupp G (2005) Quantitative evaluation of nanoindents: do we need more reliable mechanical parameters for the characterization of materials? Int J Mater Res 11: 1226-1236.
- Bower AF (2012) Mechanics of solids, Chapter 3. Electronic Text and Taylor & Francis, ISBN 9781439802472- CAT#K10131.
- Kaupp G (2009) Mechanochemistry: the varied applications of mechanical bondbreaking. Cryst Eng Comm 11: 388-403.

Select your language of interest to view the total content in your interested language

- Analytical Biochemistry Techniques
- Analytical Chemistry
- Analytical Techniques
- Antibodies Drug Conjugates
- Applied Engineering
- Bio Ceramics
- Bio inert Materials
- Bioanalysis
- Bioanalytical Techniques
- Biological Engineering
- Bionanoscience
- Biopolymers
- Brittle Materials
- Calcium Phosphate
- Capillary Electrophoresis
- Ceramic Engineering
- Ceramic Metal Oxides
- Ceramics
- Ceramics Engineering
- Clinical Biotherapeutics
- Coal Mining
- Colloid Chemistry
- Composite Materials
- Composite Materials Fabrication
- Compressive Strength
- Electronic Material Development
- Experts Opinion on Nanomedicine
- Extractive Metallurgy
- Fracture Toughness
- Geological Materials
- Human osteoblasts
- Hydrometallurgy
- Hydroxyapatite
- Implants Biology
- Implementation and Implications of Nanomedicine
- Industrial Engineering
- Injectable bone substitute material
- LC-MS principles
- Lipid based Nanoparticles
- Material Science
- Materials Chemistry
- Materials Engineering
- Materials Processing and Manufacturing
- Medicine
- Metal Casting Technology
- Metallic Materials
- Metallic Materials (Ferrous & Nonferrous)
- Metallurgical Engineering
- Metallurgy
- Mineral Processing
- Molecular Electronics
- Nano Composites
- Nano Materials
- Nano Particles
- Nano Structures
- Nanobiopharmaceuticals
- Nanobiopharmaceutics
- Nanobiotechnology
- Nanocars
- Nanocomposites
- Nanoelectronics
- Nanoemulsions
- Nanoengineering
- Nanofabrications
- Nanofluidics
- Nanohedron
- Nanoionics
- Nanolithography
- Nanomaterial
- Nanomedical Tools in Gene Therapy
- Nanomedicine
- Nanomedicine Applications
- Nanomedicine and Nanobiotechnology
- Nanomedicine in Cancer
- Nanomedicine in Drug Discovery
- Nanoparticles
- Nanoparticles and Biocompatibility
- Nanotechnology
- Nanothermite
- Nanotoxicology
- Nanotubes
- Pharmaceutical Analytical Techniques
- Polymeric Materials
- Polymeric Nanoparticle Delivery
- Porous Materials
- Prosthetic Devices
- Qualitative Analysis
- Rock Mechanics
- Semiconductors
- Solid Lipid Nanoparticles
- Spectroscopy
- Supramolecule
- Surface Mining
- TM-Joint disorders
- Therapeutic Agents
- Tissue Scaffold
- Toxicity of Nanomaterials

**10th International Conference on Emerging Materials and Nanotechnology**

July 27-29, 2017 Vancouver, Canada**10th International Conference on Advanced Materials and Processing**

August 16-17 Edinburgh, Scotland**3rd International Conference on Polymer Science and Engineering**

October 2-4, 2017 Chicago, USA**International Conference on Advanced Materials and Nanotechnology**

October 26-28, 2017 Osaka, Japan**13th International Conference and Exhibition on Materials Science and Engineering**

November 13-15, 2017 Las Vegas, Nevada, USA**14th International Conference on Functional****Energy Materials**

December 06-07, 2017 Atlanta, USA

- Total views:
**99** - [From(publication date):

June-2017 - Jul 25, 2017] - Breakdown by view type
- HTML page views :
**81** - PDF downloads :
**18**

Peer Reviewed Journals

International Conferences 2017-18